EBSD

It’s a zoo in there!

Dr. René de Kloe, Applications Specialist, EDAX

For most of us EBSD users, our day to day experience is with metals, ceramics, or perhaps rocks. For man-made materials, analysis allows us to characterise the microstructure so that we can finetune the processing or fabrication of a material for a specific application. Another common use of EBSD data is for failure analysis where the crystallographic information can be coupled to external characterisation data and deformation structures such as cracks, welds, or ductile deformation features.

Figure 1. IPF map of partially recrystallized steel (left); IQ map of quartzite rock from the Pilbara region in Australia (right).

For natural materials like rocks, the questions start to get a bit trickier as we typically do not know exactly how a rock has come to exhibit the structures that it has. In combination with other tools, EBSD can then be an invaluable tool to add crystallographic and phase information to the puzzle. This allows researchers to piece together the deformation, temperature, and pressure history of the rock. This way tiny samples can provide insight in processes on a global scale like mountain building and the motion of the continents.

A third group of materials that gets a bit less attention in EBSD analysis are biominerals, materials that are formed with a certain degree of biological control to become part of an organism. In these biomaterials, the question is not how we have produced it, or how it could be finetuned to its intended application. Here the question is how biological processes have been able to optimise a material to such a remarkable degree and the EBSD analysis is used to try to understand the biological use and control of crystallisation. Unfortunately, we rarely get to look at structures that are produced by living organisms, except possibly fossils. One of the reasons that “fresh” biomineral structures are rarely studied with EBSD is that they often contain an organic fraction that makes electron microscopy samples susceptible to beam damage. To analyse such materials, the researcher must be very careful. A single pass with the electron beam is often all you get as the structure is easily damaged. In fossilised remains of animals, the organic component has been lost or replaced by solid crystals which make its analysis somewhat easier. For example, in recent years, papers have been published on crystalline lenses in the eyes of long extinct trilobites which were formed of calcite [1] and EBSD has also been used to estimate which areas of dinosaur eggs are most likely to represent the original microstructure such that the isotope ratios from these grains can be used to estimate the crystallisation temperature of the eggs [2].

A bit closer to us is perhaps the analysis of hydroxyapatite in bones. In the SEM image this cross section of a bone consists of a fibrous framework with brighter areas containing individual hydroxyapatite grains. What is not clear from such an image is if the grain orientations in these areas are all identical or perhaps exhibit random orientation. EBSD analysis clearly shows that the apatite grains occur in small clusters with similar IPF colours or equivalent orientations, which indicates that these smaller clusters are connected in the 3rd dimension in the material.

Figure 2. BSE image cross-section of bone (left); Hydroxyapatite IPF map on a single hydroxyapatite region in bone (right).

The recent introduction of the easy recording of all EBSD patterns during a scan and performing NPAR (neighbour pattern averaging and reindexing) during EBSD post-processing have allowed dramatic improvements in the analysis of beam sensitive materials. You still have to use gentle beam currents and relatively low kV to obtain the EBSD patterns. These patterns are then very noisy and the initial maps often show poor indexing success rates, but once these have been collected you are free to find the optimum way to analyse these patterns for the best possible results. For example, beam sensitive materials like the aragonite in the nacre of shells can be successfully analysed.

Figure 3. Calcite-aragonite transition the inside of a shell: original measurement (left); after NPAR reprocessing (right).

The aragonite-calcite phase map above on the left shows the initial results of an EBSD map of the inner surface of a shell over a transition zone from the calcite “framework” on the right to the smooth nacre finish on the left of the analysis area. Directly at the interface the EBSD pattern quality is so poor that it is difficult to interpret the microstructure. The phase map on the right is after NPAR reprocessing. Now the poorly indexed zone at the transition is much narrower and the map clearly shows how the aragonite starts growing in between the calcite pillars, then forms a thin veneer on top of the calcite until it gets thick enough to create euhedral planar crystals that form the smooth nacre surface at the inside of the shell.

Figure 4. Aragonite structure from pillars to nacre: original measurement (left); after NPAR reprocessing (right).

Figure 4 shows another shell structure which is now completely composed of aragonite. In cross section the structure resembles that of the calcite pillars with the nacre platelets on top, but the initial scans do not reveal any structure in the pillars. This could be taken as evidence that the crystal structure might be damaged and cannot be characterised properly using EBSD. However, after NPAR reprocessing the crystal structure of the pillars becomes clear and a feather-like microstructure is revealed.

These fascinating biological structures don’t appear often to the average materials scientist or geologist, but if you keep an open mind for unexpected structures you can still be treated to beautiful virtual creatures in or on your samples. For example, dirt is not always just in the way. Here it poses as a micron sized ground squirrel overlooking your analysis. And this magnetite duck is just flying into view over a glassy matrix.

Figure 5. Dirt patch in the shape of a ground squirrel (left); crystal orientation map of a magnetite duck flying through glass (right).

And what to think of these creatures, a zirconia eagle that is flying over a forest of Al2O3 crystals and this micron sized dinosaur that was lurking in a granite rock from the highlands of Scotland. Perhaps we finally found an ancestor of Nessie?

Figure 6. Zirconia EDS Eagle: in zirconia -alumina ceramic (left); on PRIAS bottom image (right).

Figure 7. Ilmenite-magnetite dinosaur in a granite rock.

It is clear that “biological” EBSD can occur in many shapes and sizes. Sometimes it is literally a zoo in there!

[1] Clare Torney, Martin R. Lee and Alan W. Owen; Microstructure and growth of the lenses of schizochroal trilobite eyes. Palaeontology Volume 57, Issue 4, pages 783–799, July 2014
[2] Eagle, R. A. et al. Isotopic ordering in eggshells reflects body temperatures and suggests differing thermophysiology in two Cretaceous dinosaurs. Nat. Commun. 6:8296 doi: 10.1038/ncomms9296 (2015).

Aimless Wanderin’? – Part One

Dr. Stuart Wright, Senior Scientist, EDAX

On a recent transatlantic flight I passed the time watching one of my favorite movies: Oh Brother! There are a lot of great quotable lines in this movie. One that seems appropriate for this blog entry is from the lead character: Ulysses Everett McGill

“Say, uh, any a you boys smithies? Or, if not smithies per se, were you otherwise trained in the metallurgic arts before straitened circumstances forced you into a life of aimless wanderin’?”

Source: Rudy Wenk

While, in theory, I am “trained in the metallurgic arts”, my travels sometimes feel like “aimless wanderin’” and sometimes my mind follows suit – especially on long flights. In this series of entries for the EDAX Blog, I would like to take you on some “wanderin’s” through some of the terminology, history and personalities surrounding EBSD. Let’s begin with “texture”.

My global wanderings aren’t always aimless and I often learn some interesting things. At some recent conferences, I saw several interesting textures measured using neutron diffraction; for example, works by Heinz-Günther Brokmeier, Sven Vogel, Raul Bolmaro and others. Generally, these textures were measured over large volumes, such as from a section of a pipe, or an entire automobile component. It struck me that the use of the word “texture” has evolved to mean different things to different people.

My source of most early historical texture knowledge is Rudy Wenk. Rudy informs me that he believes the first use of the word was in an 1833 textbook by a Belgian geologist – d’Halloy to describe a directional microstructure. This seems a little ironic now as geologists tend to use the term “fabric” to describe what a metallurgist would refer to as “texture” but the evolution of these terms has also seen some wanderin’ as described in section 6 in Chapter 1 of Rudy’s 1985 book, Preferred Orientation in Deformed Metal and Rocks: An introduction to Modern Texture Analysis. I had the great fortune of learning from Rudy during a short-course on texture held at BYU when I was an undergrad as well as during his visits to Los Alamos National Lab when I was a Post-Doc. I am excited for a symposium in his honor at this year’s edition of ICOTOM in St George, Utah.

I was first introduced to the term texture in 1985 by Peter Morris, who was a visiting researcher at BYU working with Professor Brent Adams. At the time, I was employed by a Professor in the Physics Department, Dorian Hatch, to track down papers in the library (long before libraries went digital and on-line search and retrieval tools were available). I was a junior Mechanical Engineering student but had become a bit disenchanted with my coursework. I expressed to Dorian my frustration and that I was considering switching my major (Dorian was one of my leaders in our local church congregation when I was a teenager and was very helpful in offering good advice to a young university student). He recommended I go and visit with a new Professor in Mechanical Engineering named Brent Adams. When I knocked on Brent’s office door he was busy and recommended I speak with Peter. I still remember being completely lost as Peter tried to talk to me about which kind of mathematical functions would be appropriate to describe the r-dependence of the Two-Point Orientation Coherence function. Luckily, Brent popped in before I left Peter’s office completely befuddled; he brought things down a little closer to my level (if you can imagine Brent doing such a thing) and introduced me to texture. Brent was looking for someone with programming skills which I happened to have and so I joined his research team. (I got to know Peter better as part of Brent’s team particularly on a long drive from Provo, Utah to Santa Fe, New Mexico for ICOTOM 8. At one point in the drive I thought I would try out my German on Peter but was very surprised to learn that he didn’t speak German – remarkable, because if you dig out a copy of Bunge’s Texture Analysis in Materials Science you will note it was translated from German to English by Peter).

My personal introduction to texture was through the ODF or Orientation Distribution Function (another odd description as in the formal statistical sense it is actually a density function as opposed to a distribution function) per Bunge (“Zur Darstellung allgemeiner Texturen”, Zeitschrift der Metallkunde, 56, 872-874 (1965)):

“Die Orientierungsvertailung oder Textur eines polykristallinen Materials wird charakterisiert durch den Volumenateil derjenigen Kristalle, deren Orientierung zwischeng g and g + dg liegt.”

My best attempt at a translation is “the orientation distribution or texture of polycrystalline materials is characterized through the volume fraction of the constituent crystals, with orientations lying between g and dg.”

Bunge further explains in Chapter 4 of Rudy’s book entitled Preferred Orientation in Deformed Metal and Rocks: An Introduction to Modern Texture Analysis (1985):

“The texture is thus, per definition, the orientation distribution of all crystals present in the sample irrespective of their arrangement in the sample. Since the texture is defined as a statistical quantity, the sample must at least be big enough, compared to the grain size, to allow a statistically significant description. This, in turn, depends on the degree of relevance required. If we have a sample much bigger that what is required by statistical relevance, then it may be divided into volume elements V big enough to allow the statistical description of the texture. The texture can then be measured in each of these volumes elements separately. If the textures of all volume elements of the big sample are statistically identical, then the big sample is said to have a homogeneous texture. If we speak about he the texture of a material without further specification, the homogeneity is assumed. In may important cases, however, the textures of the volume elements are not the same. Such textures are called inhomogeneous, and the definition of the term “texture” become more complex (e.g., Bung, 1982c).”

In the world of EBSD, we measure textures on surfaces. We hope this is representative of the volume but oft times we know it is not. For instance, consider the following (111) pole figure measured from the surface of an aluminum sheet. It has some of the characteristics we expect for a rolled fcc material but does not exhibit the symmetry we would expect for the texture through the volume of the sheet.

(111) pole figures from two samples of rolled aluminum. Left: recent EBSD measurements on the surface of a sample. Right: X-Ray measurements from the cross-section (this pen plot is from my M.S. Thesis which formed the basis of the paper S. I. Wright and B. L. Adams (1990) An Evaluation of the Single Orientation Method for Texture Determination in Materials of Moderate Texture Strength”, Textures and Microstructures 12, 65-76.

Could the lack of symmetry be due to a lack of statistics – i.e. the volume element investigated is too small? I don’t believe so as the average grain size for this material is approximately 25 microns (always a bit tricky to estimate in deformed materials with elongated grains and with a well-defined subgrain structure) and the step size was 4µm. The scan area was 2.1 x 1.6 mm (~250,000 orientation measurements) and thus, approximately 6900 grains were sampled. In addition, the pole figure is fairly symmetric horizontally. Rather, I assume the lack of vertical symmetry in the pole figure comes from a texture gradient from the surface to the center of the sheet. So rather than calling this a texture in the classic volumetric sense it would be more correct to add “surface” as a qualifier – i.e. a surface texture.

One concern I have, is the use of the term micro-texture. I understand the point, it is the texture measured at the “micro-scale” – in the language of the quote from Bunge, a volume element at the micro-scale. But, if the area contains just a few grains, is it really a “texture”? That isn’t to say we can’t learn a lot from such measurements but, in my mind, the term texture has a statistical component to it in terms of the number of grain orientations sampled. For example, consider the following texture measurements from the same sample. Each measurement contains approximately 250,000 EBSD measurements of orientation but the step sizes are 4µm, 400nm, 40nm and 4nm. Clearly, as the sampled area becomes smaller and smaller, the measured texture becomes less and less representative of the sample as a whole. Actually, it is remarkable that the fcc rolling texture is recognizable in all but the 4nm step size. At the smallest step size, the “texture” contains just 3 grains and thus the oscillations around the major peaks arising from the spherical harmonics used to calculate the texture are relatively prominent.

(111) pole figures and orientation maps from the surface of rolled aluminum sheet from EBSD measurements at step sizes of 4µm, 400nm, 40nm and 4nm each with just over 250,000 orientation measurements.

My concern is not enough to protest the use of the word micro-texture as I think most who use the term understand the implications, but as a community we need to be aware of sampling and statistical reliability as we draw conclusions from our EBSD measurements so that our scientific wanderin’s don’t become aimless but, to quote another classic movie, “stay on target”.

(Stay tuned for some thoughts on the term “meso-texture” 😊)

My New Lab Partner

Matt Nowell, EBSD Product Manager, EDAX

It has been an exciting month here in our Draper Utah lab, as we have received and installed our new FEI Teneo FEG SEM. We are a small lab, focusing on EBSD development and applications, and without a loading dock, so timing is critical when scheduling the delivery. So, 3 months ago, we looked at the calendar to pick a day with sunshine and without snow. Luckily, we picked well.

Figure 1: Our new SEM coming off the truck.

Figure 1: Our new SEM coming off the truck.

Once we got the new instrument up and running, of course the next step was to start playing with it. This new SEM has a lot more imaging detectors than our older SEM, so I wanted to see what I could see with it. I chose a nickel superalloy turbine blade with a thermal barrier coating, as it had many phases for imaging and microanalysis. The first image I collected was with the Everhart-Thornley Detector (ETD). For each image shown, I relied on the auto contrast and brightness adjustment to optimize the image.

Figure 2: ETD image

Figure 2: ETD image

With imaging, contrast is information. The contrast in this image shows phase contrast. On the left, gamma/gamma prime contrast is visible in the Nickel superalloy while different distinct regions of the barrier coating are seen towards the right. The next image I collected was with the Area Backscatter Detector (ABS). This is a detector that is positioned under the pole piece for imaging. With this detector, I can use the entire detector, the inner annular portion of the detector, or any of three regions towards the outer perimeter of the detector.

Figure 3: ABS Detector image.

Figure 3: ABS Detector image.

I tried each of the different options, and I selected the inner annular ring portion of the detector. Each option provided similar contrast as seen in Figure 3, but I went with this based on personal preference. The contrast is like the ETD contrast is Figure 2. I also compared with the imaging options using the detector in Concentric Backscatter (CBS) mode, where 4 different concentric annular detectors are available.

Figure 4: T1 Detector (a-b mode).

Figure 4: T1 Detector (a-b mode).

My next image used the T1 detector, which to my understanding is an in-lens detector. In this mode, I selected the a – b mode, so the final image is obtained by subtracting the image from the b portion of the detector from the a portion of the detector. I selected this image because the resultant contrast is reversed from the first couple of images. Here phases that were bright are now dark, and detail within the phases is suppressed.

Figure 5: T2 Detector.

Figure 5: T2 Detector.

My final SEM image was collected with the T2 detector, another in-lens detector option. Here we see the same general phase contrast, but the contrast range is more limited and the detail within regions is again suppressed.

I have chosen to show this set of images to illustrate how different detectors, and their positioning, can generate different images from the area, and that the contrast/information obtained with each image can change. Now I have done a cursory interpretation of the image contrast, but a better understanding may come from reading the manual and knowing the effects of the imaging parameters used.

Figure 6: Always Read the Manual!

Figure 6: Always Read the Manual!

Of course, I’m an EBSD guy, so I also want to compare this to what I can get using our TEAM™ software with Hikari EBSD detectors. One unique feature we have in our software is PRIAS™, which uses the EBSD detector as an imaging system. With the default imaging mode, it subsets the phosphor screen image into 25 different ROI imaging detectors, and generates an image from each when the beam is scanned across the area of interest. Once these images are collected, they can be reviewed, mixed, added, subtracted, and colored to show the contrast of interest, similar to the SEM imaging approach described above.

The 3 most common contrasts we see with PRIAS™ are phase, orientation, and topographic. To capture these, we also have a mode where 3 pre-defined regional detectors are collected during EBSD mapping, and the resulting images available with the EBSD (and simultaneous EDS) data.

Figure 7: PRIAS™ Top Detector Image.

Figure 7: PRIAS™ Top Detector Image.

The first ROI is positioned at the top of the phosphor screen, and the resulting phase contrast is very similar to the contrast obtained with the ETD and ABS imaging modes on the SEM.

Figure 8: PRIAS™ Center Detector Image.

Figure 8: PRIAS™ Center Detector Image.

The second ROI is positioned at the center of the phosphor screen. This image shows more orientation contrast.

Figure 9: PRIAS™ Bottom Detector Image.

Figure 9: PRIAS™ Bottom Detector Image.

The third ROI is positioned at the bottom of the phosphor screen. This image shows more topographical contrast. All three of these images are complementary, both to each other but also to the different SEM images. They all give part of the total picture of the sample.

Figure 10: Defining Custom ROIs in PRIAS™.

Figure 10: Defining Custom ROIs in PRIAS™.

With PRIAS™ it is also possible to define custom ROIs. In Figure 10, 3 different ROIs have been drawn within the phosphor screen area. The 3 corresponding images are then generated, and these can be reviewed, mixed, and then selected. In this case, I selected an ROI that reversed the phase contrast, like the contrast seen with the T1 detector in Figure 4.

Figure 11: PRIAS™ Center Image with EDS Bland Map (Red-Ni, Blue – Al, Green-Zr)

Figure 12: PRIAS™ Center Image with Orientation Map (IPF Map Surface Normal Direction).

figure-12a

Of course, the PRIAS™ information can also be directly correlated with the EDS and EBSD information collected during the mapping. Figure 11 shows an RGB EDS map while Figure 12 shows an IPF orientation map (surface normal direction with the corresponding orientation key) blended with the PRIAS™ center image. Having this available adds more information (via contrast) to the total microstructural characterization package.

I look forward to using our new SEM, to develop new ideas into tools and features for our users. I imagine a few new blogs posts should come from it as well!

The Hough Transform – An Amazing Tool.

Shawn Wallace, Applications Engineer, EDAX

Part of my job is understanding and pushing the limits of each part of our systems. One of the most fundamental parts of the EBSD system is the Hough Transform. The Hough Transform role is finding the lines on an EBSD pattern. This is the first step in indexing a pattern (Fig. 1). If this step is not consistent, the quality of any indexing and any derivative data is questionable. A normal user does not really need to understand all the intricacies of every part of the system, but it still is worthwhile to understand how your data and data quality can be affected.

Figure 1: On the left are the overlaid lines found via the Hough Transform. On the right is the Indexed solution overlaid based on the Hough. The quality of the indexed solution is based on the quality of the Hough.

Figure 1: On the left are the overlaid lines found via the Hough Transform. On the right is the Indexed solution overlaid based on the Hough. The quality of the indexed solution is based on the quality of the Hough.

With that in mind, I ran an experiment on a steel sample to see how far the Hough could be pushed and still give consistent indexing. For this experiment, I used our Hikari Super at a series of different binnings between its native resolution of 640X480 Pixels at 1×1 binning down to 35×26 pixels at 18×18 binning. All pixel resolutions are noted in Table 1. I kept my Hough Settings and beam settings consistent. My only other variable was exposure to get the camera to be equally saturated at around 0.85 saturation.

I expected the lower binning Patterns to be consistent and they were (Fig. 2). All three Euler Angles between the 1×1, 2×2, 4×4, and 8×8, were within 0.4 degrees of each other. Pushing the camera and the Hough even further really surprised me though.

Figure 2: Indexed Pattern for the lower binning showed a remarkable consistency in indexing.

Figure 2: Indexed Pattern for the lower binning showed a remarkable consistency in indexing.

Figure 3: The indexing results still held their consistency even for highest binning settings used.

Figure 3: The indexing results still held their consistency even for highest binning settings used.

I expected some drop off with the consistency of the orientation when I dropped my binning to 10×10, 16×16, and even 18×18 and it did not fully materialize (Fig. 3). The range did broaden in the Euler Angles, specifically ᶲ₂’s range increased to 3 degrees, but that is change of <1% given the entire range for ᶲ₂ is 360 degrees. Table 1 shows the data is the raw form. Overall, the data is great, from low to high binning with minimal loss in in our indexing metrics (CI and Fit) and consistency in Euler Angles except for the 18×18 binning. That is where we have found our limit, specifically when it comes to indexing metrics. We see a sharp drop off in the CI. The pixilation of the pattern has gotten to a point where it is difficult to find a unique solution. This drop off is why we tell our customer that 16×16 is the limit of binning they should use for reliable, high quality data.

Table 1. Indexing Metrics and Euler Angles for all data points.

Table 1. Indexing Metrics and Euler Angles for all data points.

With all that said, most EBSD work is not on a single orientation, but a map. Does this hold true on a map? It does. In Figure 4 and Figure 5, we can see the mapping results for 2×2 binning and 10×10 binning. Both indexed at 99.9% with their average CI’s being 0.89 and 0.84 respectively, with very little change in orientations. This level of data quality across binnings is why EDAX uses the Hough. It is an amazing little tool.

Figure 4. This map was taken at 2x2 binning. Internal deformation of the grains is visible, with inclusions between relatively undeformed.

Figure 4. This map was taken at 2×2 binning. Internal deformation of the grains is visible, with inclusions between relatively undeformed.

Figure 5. This map was taken at 10x10 binning in approximately the same area as Figure 4. Again, internal deformation is showed in the larger grain, while the inclusions are undeformed.

Figure 5. This map was taken at 10×10 binning in approximately the same area as Figure 4. Again, internal deformation is showed in the larger grain, while the inclusions are undeformed.

Molecular Machines are the Future…

René Jansen, Regional Manager, Europe

The ground in the north of Holland was recently shaking and not because of an earthquake, but because Professor Ben Feringa from the University of Groningen has won the 2016 Nobel Prize in Chemistry for his work on the development of molecular machines.
Feringa discovered the molecular motor — a light-driven rotary molecular motor – which is widely recognized as a spectacular scientific breakthrough.

Electrically driven directional motion of a four-wheeled molecule on a metal surface

Electrically driven directional motion of a four-wheeled molecule on a metal surface

‘Building a moving molecule is not that difficult in itself, but being able to steer it, have control over it, is a different matter.’, he said. Years ago he already presented the first molecular motor, consisting of a molecule, part of which performed a full rotation under the influence of light and heat. He has designed many different engines since, including a molecular ‘4-wheel drive’ car. By fixating the engine molecules to a surface, he developed a nano ‘mill park’ in which the mills rotate when exposed to light. And last year he described the world’s first symmetrical molecular engine. Feringa also succeeded in putting these molecular engines to work, having them turn a glass cylinder 10,000 times their size. Amazing.

Feringa is internationally recognized as a pioneer in the field of molecular engines. One of the potential applications of his engines is the delivery of medication inside the human body.
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I recently heard an interview with him, in which he promoted the idea that universities should be playgrounds, where scientists must be able to do whatever they want to create real breakthroughs. Today, the ability of universities to create these playgrounds is limited due to a constant reduction of budgets over recent years. It would be interesting to know how the University of Groningen has managed to do this.

Another, less famous, department at the University of Groningen is working on the formation/deformation of materials which are exposed to high temperature (> 1000 degrees Celsius). Measuring EBSD patterns while temperature increases, shows that new crystals are formed at a certain temperature. Now my hopes are that this “playground” too will end up in a few years from now with a Nobel prize for a breakthrough in Materials Science.

My Turn

Dr. Stuart Wright, Senior Scientist, EDAX

One of the first scientific conferences I had the good fortune of attending was the Eighth International Conference on Textures of Materials (ICOTOM 8) held in 1987 in Santa Fe, New Mexico. I was an undergraduate student at the time and had recently joined Professor Brent Adams’ research group at Brigham Young University (BYU) in Provo, Utah. It was quite an introduction to texture analysis. Most of the talks went right over my head but the conference would affect the direction my educational and professional life would take.

Logos of the ICOTOMs I've attended

Logos of the ICOTOMs I’ve attended

Professor Adams’ research at the time was focused on orientation correlation functions. While his formulation of the equations used to describe these correlations was coming along nicely, the experimental side was quite challenging. One of my tasks for the research group was to explore using etch pits to measure orientations on a grain-by-grain basis. It was a daunting proposition for an inexperienced student. At the ICOTOM in Santa Fe, Brent happened to catch a talk by a Professor from the University of Bristol named David Dingley. David introduced the ICOTOM community to Electron Backscatter Diffraction (EBSD) in the SEM. Brent immediately saw this as a potential experimental solution to his vision for a statistical description of the spatial arrangement of grain orientations in polycrystalline microstructures.

At ICOTOMs through the years

At ICOTOMs through the years

After returning to BYU, Brent quickly went about preparing to get David to BYU to install the first EBSD system in North America. Instead of etch pits, my Master’s degree became comparing textures measured by EBSD and those measured with traditional X-Ray Pole Figures. I had the opportunity to make some of the first EBSD measurements with David’s system. From those early beginnings, Brent’s group moved to Yale University where we successfully built an automated EBSD system laying the groundwork for the commercial EBSD systems we use today.

I’ve had the good fortune to attend every ICOTOM since that one in Santa Fe over 30 years ago now. The ICOTOM community has helped germinate and incubate EBSD and continues to be a strong supporter of the technique. This is evident in the immediate rise in the number of texture studies undertaken using EBSD immediately after EBSD was introduced to the ICOTOM community.

The growth in EBSD in terms of the percentage of EBSD related papers at the ICOTOMs

The growth in EBSD in terms of the percentage of EBSD related papers at the ICOTOMs

Things have a way of coming full circle and now I am part of a group of three (with David Fullwood of BYU and my colleague Matt Nowell of EDAX) whose turn it is to host the next ICOTOM in St George Utah in November 2017. The ICOTOM meetings are held every three years and generally rotate between Europe, the Americas and Asia. At ICOTOM 18 we will be celebrating 25 years since our first papers were published using OIM.
icotom-2017
It is a humbling opportunity to pay back the texture community, in just a small measure, for the impact my friends and colleagues within this community have had both on EBSD and on me personally. It is exciting to consider what new technologies and scientific advances will be germinated by the interaction of scientists and engineers in the ICOTOM environment. All EBSD users would benefit from attending ICOTOM and I invite you all to join us next year in Utah’s southwest red rock country for ICOTOM 18! (http://event.registerat.com/site/icotom2017/)

Some of the spectacular scenery in southwest Utah (Zion National Park)

Some of the spectacular scenery in southwest Utah (Zion National Park)

With Great Data Comes Great Responsibility

Matt Nowell, EBSD Product Manager, EDAX

First, I have to acknowledge that I stole the title above from a tweet by Dr. Ben Britton (@BMatB), but I think it applies perfectly to the topic at hand. This blog post has been inspired by a few recent events around the lab. First, our data server drives suffered from multiple simultaneous hard drive failures. Nothing makes you appreciate your data more than no longer having access to it. Second, my colleague and friend Rene de Kloe wrote the preceding article in this blog, and if you haven’t had the opportunity to read it, I highly recommended it. Having been involved with EBSD sample analysis for over 20 years, I have drawers and drawers full of samples. Some of these are very clearly labeled. Some of these are not labeled, or the label has worn off, or the label has fallen off. One of these we believe is one of Rene’s missing samples, although both of us have spent time trying to find it. Some I can recognize just by looking, others need a sheet of paper with descriptions and details. Some are just sitting on my desk, either waiting for analysis or around for visual props during a talk. Here is a picture of some of these desk samples including a golf club with a sample extracted from the face, a piece of a Gibeon meteorite that has been shaped into a guitar pick, a wafer I fabricated myself in school, a rod of tin I can bend and work harden, and then hand to someone else to try, and a sample of a friction stir weld that I’ve used as a fine grained aluminum standard.

fig-1_modified
Each sample leads to data. With high speed cameras, it’s easier to collect more data in a shorter period of time. With simultaneous EDS collection, it’s more data still. With things like NPAR™, PRIAS™, HR-EBSD, and with OIM Analysis™ v8 reindexing functionality, there is also a driving force to save EBSD patterns for each scan. With 3D EBSD and in-situ heating and deformation experiments, there are multiple scans per sample. Over the years, we have archived data with Zip drives, CDs, DVDs, and portable hard drives. Fortunately, the cost for storage has dramatically decreased in the last 20+ years. I remember buying my first USB storage stick in 2003, with 256 MB of storage. Now I routinely carry around multiple TBs of data full of different examples for whatever questions might pop up.

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How do we organize this plethora of data?
Personally, I sometimes struggle with this problem. My desk and office are often a messy conglomerate of different samples, golf training aids (they help me think), papers to read, brochures to edit, and other work to do. I’m often asked if I have an example of one material or another, so there is a strong driving force to be able to find this quickly. Previously I’ve used a database we wrote internally, which was nice but required all of us to enter accurate data into the database. I also used photo management software and the batch processor in OIM Analysis™ to create a visual database of microstructures, which I could quickly review and recognize examples. Often however, I ended up needing multiple pictures to express all the information I wanted in order to use this collection.

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To help with this problem, the OIM Data Miner function was implemented into OIM Analysis™. This tool will index the data on any given hard drive, and provide a list of all the OIM scan files present. A screenshot using the Data Miner on one of my drives is shown above. The Data Miner is accessed through this icon on the OIM Analysis™ toolbar. I can see the scan name, where it is located, the date associated with the file, what phases were used, the number of points, the step size, the average confidence index, and the elements associated with any simultaneous EDS collection. From this tool, I can open a file of interest, or I can delete a file I no longer need. I can search by name, by phase, or by element, and I can display duplicate files. I have found this to be extremely useful in finding datasets, and wanted to write a little bit about it in case you may also have some use for this functionality.