13. CHARACTERIZATION OF NANO-CRYSTALLINE MATERIALS USING ELECTRON

13. CHARACTERIZATION OF NANO-CRYSTALLINE MATERIALS USINGELECTRON BACKSCATTER DIFFRACTION IN THE SCANNINGELECTRON MICROSCOPE

J. R. MICHAEL

1. INTRODUCTION

Nano-crystalline and ultra-fine-grained materials provide unique challenges in char-acterization. Until recently, the study of crystallography at the sub-micrometer scalehas been the exclusive domain of the transmission electron microscope (TEM), orif microstructural observations are not needed x-ray diffraction (XRD). Electronbackscatter diffraction (EBSD) now allows the crystallography of bulk samples to bestudied in the SEM with sub-micron spatial resolution. The obvious advantage to thestudy of crystallography in the SEM over the TEM is that the bulk specimen used inthe SEM can reduce the need for thin sample preparation and may permit larger, morerepresentative areas of the sample to be analyzed.

EBSD in the SEM has been developed for two different purposes. The oldestapplication of EBSD is for the measurement of texture on a pixel-by-pixel basis.Texture determined in this way has been termed the microtexture of the sample.The other use of EBSD is for the identification of micrometer or sub-micrometercrystalline phases. However, phase identification has not yet succeeded in addressingfeatures smaller than 100 nm [1, 2]. Both applications of EBSD add an importantnew tool to the SEM which now has the capability to study the morphology of asample through either secondary or backscattered electron imaging, the chemistrythrough energy dispersive spectrometry (EDS) or wave-length dispersive spectrometry(WDS) and the addition of EBSD permits the crystallography of the sample to bestudied.

402 II. Electron Microscopy

Important structure-property relationships may depend strongly on the crystallo-graphic texture of the sample. Properties of polycrystalline materials are determinedby the individual properties of the grains and the statistical characteristics of the poly-crystal as a whole. The information that may be obtained through EBSD are thedistributions of crystallographic orientations of the individual grains and the distribu-tion of the orientations between the grain or the grain boundary misorientations. Thedevelopment of electron backscattering diffraction (EBSD) into commercial tools forphase identification and orientation determination has provided a new insight into thelinks between microstructure, crystallography and materials physical properties thatwas not previously attainable. The combination of field emission electron sourcedscanning electron microscopes and EBSD has now allowed EBSD to be applied tonano-crystalline materials (grain size < 100 nm ) and ultra-fine grained materials (grainsize < 500 nm) [3].

Microtexture is a term that means a population of individual orientations that areusually related to some feature of the sample microstructure. A simple example ofthis is the relationship of the individual grain size to grain orientation. The conceptof microtexture may also be extended to include the misorientation between grains,often termed the mesotexture. Thus, the misorientation of all the grain boundaries ina given area may be determined. It is now possible using EBSD to collect and analyzetens of thousands of orientations per hour thus allowing excellent statistics in variousdistributions to be obtained. The ability to link texture and microstructure has lead tosignificant progress in the understanding of recrystallization, gain boundary structureand properties, grain growth and many other physical phenomena and propertiesimportant in nanocrystalline materials [4, 5].

This chapter will first describe the origin of the EBSD pattern in the SEM. Thehardware required to collect these patterns is then discussed along with the instru-mental operating conditions. The chapter will then proceed to discuss the details ofmicrotexture determination using EBSD.

2. HISTORICAL DEVELOPMENT OF EBSD

This technique has had many different names over the past 40 years. The originaldevelopers called them high angle Kikuchi patterns (HAKP) [6]. Others have usedelectron backscatter patterns (EBSP), electron backscatter diffraction (EBSD), andbackscattered electron Kikuchi diffraction (BEKP or BKD). We will use the terms elec-tron backscatter diffraction to describe this technique. The first EBSD patterns wereobserved over 40 years ago and were termed high angle Kikuchi patterns. This wasalmost ten years before the first SEM was built so a special experimental apparatushad to be constructed. The apparatus consisted of a cylindrical chamber that was linedwith a strip of photographic film. An electron beam was focused on to a tilted sampleof LiF exposing the film. This produced patterns that are every bit as good as thosecollected today on modern SEMs with modern EBSD acquisition hardware [6].

The addition of an appropriate camera system to an SEM for the observation ofEBSD patterns was demonstrated in the early 1970’s [7]. These authors demonstrated

13. Characterization of Nano-crystalline Materials using Electron Backscatter 403

that sample orientations could easily be measured from EBSD patterns and coined theterm electron backscatter patterns (EBSP). At about that same time others demon-strated some of the significant advantages held over selected area electron channelingpatterns (SAECP) [8]. Subsequently, in the late 1980’s and early 1990s, the auto-mated indexing of EBSD patterns was developed followed by systems that could scanthe sample point-by-point and determine the orientation at each point automati-cally [9]. This is the basis for determining the microtexure of a sample region. Thereare now many commercial tools that represent variations of these original develop-ments. Also in the early 1990’s, EBSD for phase identification was perfected, whichhas also served as the basis for a commercial phase identification system [10, 11]. Amore detailed review of the historical development of EBSD may be found elsewhere[12].

3. ORIGIN OF EBSD PATTERNS

EBSD patterns are obtained in the SEM by illuminating a highly tilted specimen witha stationary electron beam. Currently two mechanisms may describe the formationof EBSD patterns. In one description, the elastic scattering of previously inelasticallyscattered electrons forms the patterns. These backscattered electrons appear to originatefrom a virtual point below the surface of the specimen. Some of the backscatteredelectrons will satisfy the Bragg condition (+θ and −θ ) and are diffracted into cones ofintensity with a semi-angle of (90◦−θ ), with the cone axis normal to the diffractingplane. Since the wavelength of the electron is small, the Bragg angle is typically small,less than 2◦. These pairs of flat cones intercept the imaging plane and are imaged as twonearly straight Kikuchi lines separated by an angle of 2θ . An alternative description ofEBSD pattern formation is the single event model. In this model, it is argued that theinelastic and elastic scattering events are intimately related and may be thought of asone event. In this case the electron channels out of the sample and forms the EBSDpattern [13].

A typical electron backscatter diffraction pattern is shown in Figure 1. The patternconsists of a large number of parallel lines of intensity. These lines may appear asdiscreet lines or may appear as a band of increased intensity in the patterns. Frequently,as shown in figure 1, one of the pair of a set of Kikuchi lines will appear dark relativeto the other. This is a result of the unequal angles between the scattered beam andthe primary beam that leads to an unequal intensity transfer to one of the diffractedcones of intensity. There are many places where a number of line pairs intersect. Theintersections are zones axes and are related to specific crystallographic directions inthe crystal. The lines represent planes with in the crystal structure and the zone axesrepresent crystallographic directions within the crystal. The actual trace of the planefor a given set of diffracted lines lies exactly in between the two lines or at the centerof a band. The EBSD pattern is essentially a map of the angular relationships that existwithin the crystal.

The spacing between a pair of lines is inversely related to the atomic spacing ofthose planes within the crystal. The EBSD pattern is an angular map of the crystal


Figure 1. A typical EBSD pattern from Fe2O3 (hematite, rhombohedral) recorded at 20 kV with aCCD-based slow scan camera.

so we should relate the distance between any pair of lines to an angle. The angle fordiffraction from a given set of planes is called the Bragg angle and is given by:

λ = 2d sin θ (1)

Where d is the spacing of the atomic planes, λ is the wavelength of the electron andθ is the Bragg angle for diffraction. From this equation, it can be clearly seen that asthe interatomic spacing increases the Bragg angle must decrease. Thus, smaller atomicplane spacings will result in wider line pairs.

3.1. Collection of EBSD Patterns

Acquisition of EBSD patterns in the SEM is relatively easy. Some of the earliest imagesof these patterns were obtained by directly exposing photographic emulsions insidethe specimen chamber of an SEM [6]. Although this is still an option, all commercialsystems now use a phosphor screen inside the sample chamber of the SEM that isimaged by an external camera through a leaded glass window. A schematic of a typicalexperimental setup is shown in Figure 2. The tilt angle of the sample with respect tothe normal to the electron beam is typically set to about 70◦. The exact value is notcritical, as patterns have been obtained from tilt angles as small as 40◦. Figure 3 showsan actual SEM chamber fitted with an EBSD camera.

There are no specific requirements on the electron beam parameters. The onlyrequirement is that the diffracted electrons reaching the recording medium (either aphosphor screen or a photographic emulsion) must have a sufficient kinetic energy


Tilted Sample

Diffracting planes

Electron Beam

diffracted cone of electrons

Figure 2. Typical EBSD experimental configuration. Note the relationship between the electron beam,the tilted sample and the phosphor screen.

to produce an appropriate response (generation of light) in the recording mediumand must have a large enough current to produce a measurable signal. Generally, thismeans that a beam current of greater than 0.1 nA and an accelerating voltage of greaterthan 2.5 kV must be used. More typical values are an accelerating voltage of 10 kVand a beam current of 0.5 nA or larger. Due to the response of the commonly usedphosphor medium, it is usually not possible to work at very low beam currents andlow accelerating voltages.

The detection of the patterns has been accomplished using many different types ofcameras and photographic film. The most commonly used camera is either a TV ratelow light level camera or a CCD-based imaging system. The CCD-based systems aremost versatile as a variety of parameters, such as binning and gain levels, can be setto provide excellent EBSD patterns over a range of microscope operating conditions.Figure 4 shows a series of EBSD patterns collected from Ni on a CCD-based camerasystem using a range of accelerating voltages.

3.2. Automated Orientation Mapping

The study of the orientation of individual grains within a polycrystalline assemblagehas become indispensable in materials science. It has only been recently that EBSDin the SEM and automated pattern indexing and orientation measurement have been


Figure 3. Arrangement of the EBSD camera (arrowed) with respect to the specimen in a SEM.

combined. A use of EBSD in the SEM is to determine the local relationships betweenthe microstructure of the sample and crystallography. This may simply involve thedirect correlation between the orientations of the grains in a polycrystalline material inwhich some interesting event is occurring or determining the orientation relationshipbetween some precipitate and the matrix phase or the relationship between the fracturepath and crystallography. These cases require only a few measurements at discreetpoints and do not need fully automatic orientation measurement, although automationmakes this work quite easy. Other cases, for example, the measurement of texture,require many thousands of measurements so that the distribution of orientations maybe obtained.

There are two ways that automated orientation mapping can be accomplished.Stage scanning requires the electron beam to be held stationary and the sample israstered by moving the sample stage. Beam scanning utilizes the scan coils in the SEMcolumn to scan the beam over a stationary sample. In the early days of EBSD, thepreferred method of scanning the beam over the sample surface was stage scanning.Stage scanning has a number of important advantages. The first advantage is that there


Figure 4. EBSD patterns of Ni collected at a) 5 kv, b) 10 kV, c) 15 kV and d) 20 kV e) 25 kV andf ) 30 kV.

is no defocus associated with moving the beam over the highly tilted sample. This isonly possible if the x and y motor drives for the stage are in the plane of the tiltedsample. The second advantage is related in that the beam position with respect to theEBSD phosphor screen is fixed and thus the calibration (the pattern center and thedistance between the sample and the phosphor screen) does not vary. The disadvantageof stage scanning is that it is much slower than beam scanning. Generally, the fastest


stage scan requires about 1 second per pixel. The other disadvantage is that the accuracyof movement is not much better than 1 μm and is therefore not of much use to thestudy of nano-crystalline materials.

Beam scanning is the more commonly used technique for automated orientationmapping. For beam scanning, the scan generator of the SEM is disabled and the elec-tron beam is rastered by an external scan generator controlled by the EBSD computersystem. The advantage of this system is that the time required to reposition the beamis quite short when compared to the time required to collect and index the EBSDpattern. Thus, beam scanning is much more rapid when compared to stage scanning.The disadvantages of stage scanning are related to the motion of the electron beam withrespect to the SEM column axis and the EBSD phosphor screen. The tilted sampleresults in an out of focus condition at the highest and lowest positions on the specimen.This can be partially compensated for by using dynamic focus corrections and smallobjective apertures, but the defocus at low magnifications cannot be eliminated. Theother related disadvantage is that the calibration of the EBSD system changes withthe position of the beam on the sample. This effect may be reduced by automaticallycorrecting the calibration of the system as a function of beam position on the sample.These difficulties are primarily noted at lower magnification where there may also besome distortion of the microscope scans anyway. At higher magnification these disad-vantages are not significant and beam scanning is recommended over stage scanningfor the study of fine-grained or nano-crystalline materials. One very useful mode ofoperation is a mix of beam scanning and stage scanning. In this mode the beam isscanned over an area of interest. After the scan is complete the stage is moved and thenanother beam scan is commenced. Accurate positioning of the areas allows the mapsto be knitted together to produce larger area maps.

At each pixel the EBSD pattern is acquired using a low-light level video or CCD-based camera that images the phosphor screen. The bands or lines in the pattern arethen detected through the use of the Hough or Radon transform of the image [14].Once the lines are found, the angles between the bands can be calculated and comparedto a look up table of angles that is constructed from the crystallography of the sample.Once a consistent set of indices can be assigned to the bands, the orientation of thepixel can be described with respect to some external set of reference axes. This processis then repeated at every pixel throughout the area of interest [15].

4. RESOLUTION OF EBSD

4.1. Lateral Resolution

In order to use EBSD for the study of nano-crystalline materials we must under-stand the spatial resolution of the measurement. The spatial resolution of EBSD isstrongly influenced by the atomic number of the material to be studied, the accel-erating voltage of the SEM, the focused probe size and the sample tilt. For EBSDto be useful for nano-crystalline materials, each of these parameters must be care-fully set to achieve high spatial resolution required for the study of nano-crystallinematerials.


Figure 5. Resolution differences resulting from the use of a highly tilted sample. Note that boundariesparallel to the tilt axis are sharper than those perpendicular to the tilt axis. a) Orientation map with respectto an in-plane direction. Pixels not indexed are shown in black, indexed pixels in color. b) Pattern qualityimage of same area. (See color plate 7.)

The spatial resolution of EBSD is strongly influenced by the sample tilt. The highsample tilts required for EBSD results in an asymmetric spatial resolution parallel andperpendicular to the axis of tilt. This asymmetric spatial resolution also influences thedepth from which the EBSD pattern is generated. The resolution parallel to the tiltaxis is much better than the resolution perpendicular to the tilt axis due to the highsample tilt angles used to acquire EBSD patterns. The resolution perpendicular to thetilt axis is related to the resolution parallel to the tilt axis by:

Lperp = Lpara(1/cos θ ) (2)

Where θ is the sample tilt with respect to the horizontal, Lperp is the resolution perpen-dicular to the tilt axis and Lpara is the resolution parallel to the tilt axis. The resolutionperpendicular to the tilt axis is roughly three times the resolution parallel to the tiltaxis for a tilt angle of 70◦ and increases to 5.75 times for a tilt angle of 80◦. Thus,it is best to work at the lowest sample tilt angles possible consistent with obtaininggood EBSD patterns. The difference in resolution parallel to and perpendicular to thetilt axis is apparent in Figure 5. The image in this figure is generated by measuringthe pattern sharpness or quality at each pixel within the map, and then displayed asa gray scale. The pattern sharpness or contrast is reduced when the electron beam isvery close to a grain boundary due to contributions to the patterns from the grainson either side of the boundary resulting in an EBSD pattern containing overlappingpatterns. The grain boundary parallel to the tilt axis is wider than the boundary per-pendicular to the tilt axis demonstrating the reduced resolution perpendicular to the tiltaxis.

The spatial resolution of EBSD is directly related to the electron beam size orfocused probe size at the sample. Smaller beam sizes result in higher spatial resolution


0 5 10 15 20

#b

acks

catt

ered

ele

ctro

ns

Energy (keV)

Figure 6. Energy distribution of backscattered electrons for normal incidence and a 70◦ sample tilt.

for EBSD measurements. For this reason, when studying nano-crystalline materials,an SEM with a field emission source must be considered mandatory for producingEBSD patterns. Although EBSD is a technique that utilizes backscattered electrons,the resolution of the technique is generally much better than can be achieved withstandard backscattered electron imaging in the SEM. This is a direct result of theway in which the patterns are formed. Monte Carlo electron trajectory simulationsare used to study the energy distribution of the electrons that exit the sample and anexample of this is shown in Figure 6, where the energy distribution of the backscatteredelectrons is plotted for an untilted sample and one tilted 70◦ from the horizontalposition. This plot is drawn for an initial electron beam energy of 20 kV interactingwith a Ni specimen. The number of initial trajectories was the same for each sampleorientation. There are two important points to note from this figure. First, it is veryclear that tilting the sample to a high angle with respect to the normal results in amuch higher backscatter yield. This is important as the increased signal improves thequality of the EBSD patterns and decreases the acquisition time. The other importantpoint is the shape of the backscattered electron energy distribution. It is apparent fromFigure 6 that many of the backscattered electrons from the tilted sample have nearly theinitial electron beam energy. The large peak in the tilted sample backscattered energydistribution is only a few hundred eV less than the initial beam energy. The electronsin the sharp energy peak contribute to the crystallographic information in the EBSDpattern, while the electrons in the remainder of the distribution contribute to the


Figure 7. Comparison of EBSD resolution with large and small electron beam sizes. a) Pattern qualitymap for a large electron beam size, b) pattern quality map for a much smaller electron beam size, c)orientation map corresponding to beam size used in a and d) orientation map corresponding to beam sizeused in b. Black pixels in the orientation maps are where the EBSD patterns could not be indexed. Arrowsindicate the same areas in each image. All maps are 120 × 80 pixels with a step size of 100 nm and werecollected in 10 minutes. (See color plate 8.)

overall background intensity of the EBSD pattern. The sample that is normal to theelectron beam has a broad spectrum of energies associated with the backscatteredelectrons which mainly contribute to producing the background signal in an EBSDpattern with no visible crystallographic information. Under certain special conditionsit has been shown that EBSD patterns may be generated at normal incidence, but thereare significant challenges to utilizing this geometry [16].

Monte Carlo electron trajectory simulations have shown that the low loss backscat-tered electrons emerge from the specimen surface very near the initial beam impactpoint [17]. This causes the resolution of EBSD to depend mainly on the electronprobe size. The spatial resolution is also a function of the accelerating voltage and theatomic number of the sample. Lower accelerating voltages will improve the spatialresolution due to the decreased range of the backscattered electrons within the sam-ple as long as the electron beam size does not increase (may be difficult to achievein practice). Higher atomic number samples will have improved spatial resolutionwhen compared to lower atomic number specimens. This is a result in the decreasedrange of the backscattered electrons in the higher atomic number sample. Figure 7can be used to illustrate the effect of electron beam size. Figure 7a and b are grayscale maps of the EBSD pattern sharpness at each pixel. Figure 7a was acquired using


Table 1. Lateral Spatial Resolution in nm for Al andCu vs. SEM Operating Voltage

10 kV 20 kV 30 kV

Al 60 nm 200 nm 300 nmCu 30 nm 50 nm 80 nm

non-optimum beam parameters to produce a rather larger electron beam size. Note thelack of sharpness of the grain boundaries and the overall low image quality. Figure 7bwas generated using parameters that provided a much smaller beam diameter than in 7awith adequate beam current. Note the obvious blurring of the grain boundary regionsin the patterns quality map of Figure 7a. The blurring of the image near the grainboundaries is due to overlap of the beam onto two grains and results in two superposedpatterns. Figure 7c and 7d are orientation maps of the same region. Black is used toshow positions on the sample from which the EBSD patterns could not be indexed. Itis immediately obvious that the smaller electron beam size used in Figure 7b and 7d hasresulted in a higher percentage of indexed patterns. This demonstrates the importanceof utilizing the smallest electron beam size that has sufficient current to generate EBSDpatterns in fine-grained or nano-crystalline materials.

There have been some studies that have measured the spatial resolution of EBSD.We will discuss only those studies that have used field emission electron sourced SEMinstruments. Generally, the spatial resolution of the technique has been determined byplacing the electron probe near a grain boundary and noting the minimum distancefrom the grain boundary that allows EBSD patterns to be obtained with no overlapfrom the opposite grain. This technique has produced results that vary, but can beused as a guideline for the use of EBSD in fine-grained materials. One study of Nifound that the lateral spatial resolution was 150 nm, 100 nm and 50 nm at 30 kV,20 kV and 10 kV respectively [18]. Another study has shown that for Al the lateralspatial resolution of about 20 nm [19]. A more comprehensive study that utilizedexperimental measurements and electron trajectory simulations has determined theimproved spatial resolutions resulting from lower accelerating voltage SEM conditions,as shown in Table 1 [20]. Table 1 shows that the spatial resolution of the technique isquite high and suitable for studying nano-crystalline materials and can be better than50 nm. It is clear that the best spatial resolutions are achieved by operating at loweraccelerating voltages and in higher atomic number materials. A purely experimentalstudy of EBSD spatial resolution has shown a resolution of about 20 nm in Al and asgood as 9 nm in brass. These results are calculated based on the fraction of indexedEBSD patterns as a function of grain size and not on a physical measurement of theEBSD interaction volume [21].

4.2. Depth Resolution

The depth resolution of EBSD is also a strong function of the specimen atomic numberand the operating voltage of the SEM and is more difficult to measure experimentally.The maximum distance that a Bloch wave can travel unscattered is about 2 to 3


extinction distances for the given reflection. Thus the depth resolution has beendescribed as 2 or 3 times the many-beam extinction distance for the acceleratingvoltage and the material [22]. This leads to a worst-case depth resolution at 20 kVof about 100 nm for Al and about 50 nm for Ni. As in the lateral resolution, thedepth resolution is improved by operating at lower operating voltages or for higheratomic number samples. Monte Carlo electron trajectory simulations have been usedto predict the depth resolution of EBSD based on the use of a 90% energy loss cutoff with a resulting depth resolution at 20 kV of 100 nm for Al and 20–30 nm for Ni[20]. These results represent a worst-case value due to the generous energy cut off.In general the depth resolution appears to be similar to the lateral resolution discussedpreviously. An example of thin film analysis is discussed in the section 6.1.

5. SAMPLE PREPARATION OF NANO-MATERIALS FOR EBSD

The main sample preparation requirement for EBSD studies is that the sample sur-face shall be clean, representative of the bulk of the material and free from damage ordeformation resulting from the preparation process. Some materials, like epitaxial orheteroepitaxial layers may require no additional sample preparation steps and EBSDpatterns may be obtained from the sample in the as-deposited condition. Other sam-ples may require more complex procedures to produce good samples for EBSD. Fornano-crytalline materials it is imperative that proper sample preparation procedures arefollowed as the grain size of the material approaches the resolution of the EBSD tech-nique. Improper sample preparation can result in degraded patterns or the total loss ofthe EBSD patterns due to surface deformation or damage induced by the preparationtechnique.

For larger samples of nano-crystalline materials, it may be possible to employ con-ventional specimen preparation techniques. Standard metallographic mounting andpolishing can produce samples suitable for EBSD. In some cases it may be neces-sary to follow the mechanical polishing with a chemical attack etch to remove thelast remnants of damaged material caused by the mechanical polishing. This tech-nique has been shown to work well for many ceramics as well as metal samples.One of the best techniques for EBSD preparation is electropolishing, similar to thetechniques used for TEM sample preparation. Unfortunately, the small sample sizeoften associated with nano-crystalline materials, makes electropolishing difficult orimpossible.

The small size of many nano-crystalline samples can make conventional samplepreparation difficult or impossible. One very useful tool for the preparation of nano-crystalline materials for EBSD is the focused ion beam (FIB) tool. FIB tools have beenused for many years to produce thin samples for transmission electron microscopy(TEM) where the ion beam is used to micromachine a thin membrane from thesample. The advantage to FIB preparation is that the sample can be prepared froma very specific location. Once the thin sample is prepared using FIB the sample ismounted on a suitable substrate, usually a carbon coated TEM grid using the sameprocedures used for TEM sample preparation [23]. This method does not allow any


Figure 8. Combined EBSD and TEM of a FIB prepared cross section through a fatigue crack in a NiMEMS device. a) FIB cross section that shows the location of the section with respect to the fatigue crack,b) EBSD map of area around fatigue crack, c) annular dark field STEM image of same sample as in b.Arrows indicate the same point in all three images. The orientation map is 300 × 200 pixels with a stepsize of 25 nm and required 30 minutes to collect. (See color plate 9.)

subsequent treatment of the sample. Alternatively, the sample can be attached to theside of the grid using the metal deposition capability of the FIB so that subsequentlow energy ion milling may be performed [24]. The main disadvantage to FIB samplepreparation is the limited sample size (about 50 μm wide by 20 μm deep). Thus, thearea or feature of interest must be small and near the surface of the sample, althoughthis is not often a problem with nano-materials. Even with the sample size restriction,FIB has been used to prepare useful samples for EBSD analysis [25, 26].

Figure 8 shows an example of a FIB section prepared from the vicinity of a crackin a fatigued electrodeposited Ni sample. The sample was mounted onto a carbonfilm and an orientation map was obtained with no further preparation. We havefound this to work for many transition metals and many ceramic materials. Othermaterials, like Si and compound semiconductors, require addition preparation dueto damage to the crystal structure caused by the ion beam. These materials requirefurther low energy ion polishing in order to obtain good EBSD patterns. Low voltage(2–3 kV), low-angle (3–4◦) argon ion polishing following FIB micromachining, for


a very short time is all that is needed to produce good EBSD patterns from thesematerials.

One additional advantage of the FIB technique is that the sample may be madesufficiently thin for TEM analysis that could follow the EBSD analysis. Thus, correla-tion of the EBSD orientation maps with TEM images of the material is relatively easyand can be a powerful route to nano-materials characterization. An example of this isshown in Figure 8b and 8c, which are an orientation map of the fatigue crack regionof electrodeposited Ni sample and a STEM annular dark field image of the same area.This ability to perform correlative imaging is very powerful and allows the sample tobe more fully characterized. There are a few texts that discuss specimen preparationfor EBSD, but additional sample preparation details have been published [27].

6. APPLICATIONS OF EBSD TO NANO-MATERIALS

6.1. Heteroepitaxy of Boron Arsenide on [0001] 6H-SiC

Icosohedral boron arsenide (B12 As2, rhombohedral, hexagonal parameters, a =0.615 nm, c = 1.191 nm) is a wide-band gap semiconductor exhibiting exceptionradiation hardness and is potentially useful for the fabrication of beta-cells for directnuclear to electrical energy conversion. Chemical vapor deposition has been shownto produce crystalline films of As2B12 on 6H-SiC [28]. Characterization of the filmsis important to understand the properties and the performance of energy conversiondevices. Figure 9a is a SEM image of the surface of the boron arsenide film. The filmsare quite smooth and were shown to have useful electrical properties. Figure 9b and9c are orientation maps of a similar area. The thin black lines superposed on bothimages in 9b and 9c represent high angle grain boundaries. Figure 9b is a map that iscolored with respect to an in-plane direction and shows that two rotational variantsare present. It is important to note that when utilizing orientation maps one mustusually prepare maps that represent orientations in two orthogonal directions to get acomplete picture of the sample texture as these are equivalent to inverse pole figuresand suffer the same disadvantages. Figure 9c is an orientation map that is coloredwith respect to the surface normal. The red color over the entire sample demonstratesthat all the rotational variants have the same crystallographic direction parallel to thesample surface or the film growth direction. Figure 9b and 9c show that all areas ofthe film have grown with [0001]B2As12//[0001]6H-SiC as shown by the red colorindicating that the film has a [0001] plane parallel to the surface. Figure 10 a and b,are inverse pole figures from the map shown in Figure 9b and c. Figure 10a is thein-plane inverse pole figure that clearly shows the two rotational domains consisting of60◦ rotations about the <0001> axis as shown in the inverse pole figure of 10b. Brightand dark-field TEM images are shown of the boron arsenide film in Figure 11a andb. Note that the film is 100 nm thick and as shown in the bright field images consistsof domains of different rotational orientations as well as translational variants. Thereare also many smaller translational variants that are not detected using using EBSDmapping, making TEM imaging an important component of understanding these thinfilms [29].


Figure 9. Heteroepitaxy of Boron arsenide thin-film deposited on 6H-SiC. a)SEM image of sampletilted for orientation mapping, b) Orientation map of the in-plane texture, c) Orientation map of theout-of-plane, or surface normal orientation. Lines on the orientation maps represent boundaries betweenrotational variants. The map is 450 × 350 pixels with a step size of 100 nm and was collected in 6 hours.(See color plate 10.)

6.2. Electrodeposited Ni for MEMS Applications

Electodeposition of metals has now become quite common process in many areas ofmicro-electronics and micro-electro mechanical systems (MEMS). Metals are neededfor some MEMS devices due to the need to have thicker structures that are capableof transmitting large sustained loads or torques [30, 31]. The microstructural scale ofelectrodeposited metals can be quite small. In order to optimize the process and tounderstand the resulting properties of the electrodeposit it is important to understandthe crystallographic texture of electrodeposited metals for MEMS applications. Thereare two examples shown in this section, electrodeposition of Ni onto an Au seed layerand the electrodeposition of Ni for the fabrication of small mechanical parts using theLIGA process.

EBSD orientation studies were carried out in order to understand the electro-decomposition of Ni over patterned lines on a Au seed layer. Due to the small size ofthe features involved it was necessary to use the FIB to prepare cross sections throughthe electrodeposits. These sections were mounted on Cu TEM grids and then mountedin the SEM for orientation mapping. Figure 12 shows an overall view of one of theelectrodeposited Ni lines. Figure 12a is a pattern quality map of the sample area. These


a b

Figure 10. Inverse pole figures formed from the data shown in Figure 9b. a) Inverse pole figure withrespect to the x direction that shows the two 60◦ rotational variants b) Inverse pole figure with respect tothe surface normal demonstrating the slight misalignment of the basal planes from each rotational variant.(See color plate 11.)

Figure 11. TEM cross sectional images of the boron arsenide films of Figure 9. a) Bright field TEMimage that shows the large number of small translational domains in the sample that are not detected byEBSD, b) Dark field TEM that shows one of the rotational domains in contrast. Surface layers are due toFIB deposited Pt to aid in sample preparation.


Figure 12. Ni electrodeposited lines on patterned substrate. The thin layer at the base of the deposit is athin Au layer. a) Pattern quality map. b) Orientation map with respect to the deposition direction. c)Stereographic triangle color legend for the orientation map in b. The map is 175 × 175 pixels with a stepsize of 25 nm and was collected in 50 minutes. (See color plate 12.)

images provide excellent detailed information about the general microstructure of thesample. The microstructural details are clearly visible in the pattern quality image.Figure 12b is an orientation map with respect to the growth direction. This map iscolored with respect to the stereographic triangle shown in Figure 12c. The narrowband of blue colored grains (indicated with the arrow) at the bottom of the image isthe Au seed layer. These grains have formed with a <111> fiber texture. Figure 13 isa higher resolution image that shows the details of the electrodeposit at the Au seedlayer. Figure 13a is a pattern quality map that shows the microstructure of the deposit.Figure 13b is an orientation map with respect to the growth direction of the Ni. Itis apparent that the Au seed layer has formed with a <111> fiber texture (refer tothe color legend in Figure 12c) while the Ni layer has grown with a more randomtexture. The largest of the Au grains is about 100 nm in diameter demonstrating thatthe resolution of the technique is quite good. From this work, the growth texture ofthe Ni deposit is not strongly influenced by the texture of the Au seed layer.


Figure 13. Higher magnification of the electrodeposit shown in Figure 12. a) Pattern quality map, b)Orientation map with respect to the growth direction using the color legend shown in 12c. The fiberoriented Au is at the bottom of the image. The map is 200 × 200 pixels with a step size of 10 nm and wascollected in 1 hour. (See color plate 13.)

LIGA, an acronym of the German words “Lithographie, Galvanoformung, Abfor-mung,” is a microfabrication process in which structural material is electrodepositedinto a photolithographical realized mold. LIGA permits the fabrication of parts withextremely small cross sections, on the order of a few micrometers with accuracy bet-ter than one micrometer [30]. The microstructure of LIGA formed parts dependsprimarily on the processing parameters during electrodeposition of the metal. Thetexture of the deposit is important in determining mechanical properties. Thus carefulcharacterization and control of LIGA materials and processes is required to producehigh quality, reliable components. The small physical size of many LIGA parts makescharacterization by standard x-ray techniques difficult or impossible. EBSD orientationmapping provides the needed resolution and the ability to sample a larger area thanTEM analysis makes EBSD orientation mapping ideal for the study of these materials.

X-ray diffraction of these materials has been used to determine their macro-texture, but this does not give any indication of the relationship between textureand microstructure. EBSD orientation mapping was carried out to determine the rela-tionship between the texture of an electroformed deposit and the grain size and shape.Also, the distributions of grain boundary orientations was of interest. For this analysisthe sample was mounted and polished using standard metallographic procedures sothat the deposit could be studied in cross section. Figure 14a is a pattern quality mapof an NiMn electroformed part that shows there is a duplex grain size with somelarge columnar grains separated by smaller areas of equiaxed grains. Figure 14b is anorientation map of the area colored with respect to the growth direction. Refer to thestereographic triangle in Figure 12c for the color legend. It is immediately obvious thatthe long columnar grains have a <110> direction parallel to the growth direction and


Figure 14. Electroformed NiMn for LIGA MEMS devices. a) Pattern quality map that shows the grainstructure of the deposit b) Orientation map colored as shown in Figure 12c. The large green shadedgrains represent a <110> fiber texture. Smaller grains are randomly oriented with respect to thegrowth direction. The map is 200 × 600 pixels with a step size of 20 nm and was collected in 2.5 hours.(See color plate 14.)

the smaller equiaxed regions are more randomly oriented. From this data the volumefraction of <110> fiber oriented grains can be determined by selecting only thosepixels with orientations within 10◦ of a <110> fiber texture demonstrating that about50% of the grains have a <110> fiber orientation with respect to the growth direction.This is important for a full understanding of the mechanical properties of this material.

There have been a number of recent studies that have shown that the distributionof grain boundary misorientations can influence the performance of polycrystallinematerials. Early studies focused on the corrosion properties of a variety of metals[32]. Further studies have shown that by increasing the fraction of special boundaries(typically defined as grain boundaries near the low � misorientations) the corrosion


Figure 15. Representations of the distributions of the grain boundary misorientations in the NiMnelectrodeposits. a) Grain boundary misorientation distributions, solid line is experimental data from mapshown in Figure 12, dotted line is the distribution for a random arrangement of grains. b) Pattern qualitymap with twin boundaries shown in white.

behavior, creep resistance and weldability of the material can be enhanced [33]. Oneimportant concern with the electrodeposited NiMn material was the influence ofgrain boundary misorientations on the annealing behavior of the material. Previouswork in Ni deposits has shown that the material recrystallized in a manner that max-imized the number of low energy (twin type or �3 or a 60◦ rotation about the<111>) grain boundaries [31]. Therefore, it is important to know the misorientationdistribution of the as-deposited material. Figure 15a shows the grain boundary mis-orientation distribution for the NiMn sample. This information is calculated from theorientation data so no additional scanning needs to be performed. The data from theNiMn electrodeposit is compared with the expected distribution of grain boundarymisorientations derived ffrom the random orientations of cubes [34]. It is clear thatthere is a large difference between the expected random distribution and the measureddistribution. The most striking difference is the large number of 60◦ misorientations.These rotations can be shown to be about the <111> and are therefore �3 twintype boundaries. Figure 15b is a pattern quality image with the twin type bound-aries shown in white. The high density-of-twins in the microstructure are importantto understanding the recrystallization behavior and the response to high temperatureexposures of electrodeposited Ni.


Figure 16. shows a FIB prepared cross section through an unreleased multilayer polycrystalline Si MEMSdevice. The sample was micromachined using the FIB in a manner similar to that used to make TEMsamples and then attached to the grid using Pt deposition.

6.3. Polycrystalline Si For MEMS Applications

Polycrystalline silicon thin films are an important component of many advanced MEMSstructures. Reliability of these devices is closely related to the mechanical performanceof the thin films of polycrystalline silicon used to fabricate the devices. The needfor an understanding of these structure-property relationships in these films is mademore important by the fact that the microstructure may consist of highly texturedcolumnar grains [35, 36]. Many polycrystalline silicon MEMS devices are producedby a surface micromachining that uses silicon oxide coating as sacrificial spacers betweenthe structural layers of polysilicon.

Characterization of these types of materials are complicated by the small size of thedevices which in many cases may have features that do not exceed 1 μm. EBSD is idealfor determining the grain structure and texture of the multiple layers of polysiliconused to produce these devices. Preparation of samples is best performed with the FIB.In the case of silicon and the results shown here, subsequent low energy ion millingis required to produce EBSD patterns of sufficient quality for automated mapping.


Figure 17. Polycrystalline structures for MEMS devices. a) Pattern quality map, the single crystal Sisubstrate is at the bottom of the image. The polycrystalline Si layers are separated by oxide. b) Orientationmap with respect to the growth direction. The map is 600 × 500 pixels with a step size of 25 nm and wascollected in 8.5 hours. (See color plate 15.)

Figure 15 shows a FIB prepared cross section through an unreleased (the oxide spacershave not been removed) multilayer MEMS device. The sample was micromachinedusing the FIB in a manner similar to that used to make TEM samples [23]. The crosssection was attached to a support grid by depositing Pt at the sample grid junction,due to the need for subsequent low energy ion milling to allow EBSD patterns to beacquired. This arrangement is robust and allows the sample to be handled with littledanger of damage. The sample was low-energy argon ion milled for 30 seconds at 3 keV.After ion milling excellent EBSD patterns were obtained from the silicon. Figure 16ais a pattern quality map of the multilayer device. Note that the individual levels ofpolycrystalline silicon appear to be made up of multiple layers. Figure 16b is a orien-tation map with respect to the growth direction. Note that the substrate silicon is redindicating a <100> direction normal to the sample surface (please refer to Figure 12cfor the complete color legend). Each layer has a slightly different microstructure due tothe accumulated time at temperature for the first layer as compared to the subsequentlayers. The layer furthest from the substrate is more columnar when compared to thefirst layer. In addition there are slight textural differences from the first layer to thelast. The first layer represents annealed material and exhibits a tendency for a <111>


fiber texture as compared to the last layer which has not been at temperature for aslong and exhibits a tendency to a <110> fiber. These differences in texture resultin different mechanical properties for each layer, which assists in the proper design ofMEMS components.

7. SUMMARY

EBSD should now be considered a standard analytical accessory for the SEM. The spa-tial resolution of EBSD is quite high (as high as a few nm, depending upon material)and is suitable for the study of many nano-crystalline materials. For applications ofEBSD where the highest spatial resolution is required, an SEM with a field emis-sion electron source is mandatory as the resolution of the technique is most directlydetermined by the electron beam size. Further developments the electron optics ofthe SEM should allow even higher spatial resolutions. For nano-crystalline materi-als, EBSD is most useful for the mapping of grain orientations and grain boundarymisorientations as demonstrated in the examples shown in this chapter.

ACKNOWLEDGEMENTS

Special thanks to my colleagues at Sandia National Laboratories for bring their specialmaterials problems to my attention. In particular, I would like to thank Terry Aselagefor the boron arsenide work, Sean Hearne for the nucleation of electrodeposited Nistudy, Brad Boyce for the fatigue samples and Tom Buchheit and Steve Goods for theLIGA NiMn material.

Sandia is a multiprogram laboratory operated by Sandia Corporation, a Lock-heed Martin Company, for the United States Department of Energy under contractDE-AC04-94AL85000.

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13. CHARACTERIZATION OF NANO-CRYSTALLINE MATERIALS USING ELECTRON