SFCM 11/12 23: ELECTRON CHANNELING AND DIFFRACTION FOR ORIENTATION ANALYSIS AND DEFECT IMAGING

dias a todos tenemos una ID Fiona special de los seminary owes Fronteras con una botado estados unidos el professor martin creamed michigan state university el esta aqui invitado por los servicios de negocio nado cattiva esta dando clases en la universidad lo que pasa es que buen ademas lame-o’s pedido que nose que NOS quiere una conferencia de investigación LS un experto reconocido en análisis de feo y n Fratto graffia e tamanna trabajado muchisimo mi Claro’s copia de metales que para nosotros a son team avital y le hemos podido que NOS hablan poco de souspennier special day electron channel Martin welcome to the material science department thank you for coming thank you for this conference and please it is guitar yeah sorry es la que hay oh yeah swing ok fur so I want to thank you very much for hosting me here and please be able to participate in the 16 program we’ve had some of your students work with us and I guess one of our students is over here right now and I hope this can be able lasting relationship with Michigan State I also apologize for my very poor Spanish I’ve said about all the Spanish I know so far so it very sorry for that but yes we have to do that yesterday actually gave a talk over in David Morris’s lab at what national metallurgical research center I don’t know the exact translation by talk some there about a nice at tropic classic behavior in amateur nucleation in in titanium and today I’m going to talk about something a little bit different I’m going to overlap a little bit with that talk but I can talk really more about a technique in some techniques for examining plasticity that are something that I’ve been working on for now about 10 12 years or so that allow us to do things that we’ve traditionally done with transmission electron microscopy maybe do some of those sorts of studies using scanning electron microscopy which makes things a little bit easier and so I wanted to introduce you to those ideas obviously we want to thank the people who supported this work over the years I think like in your country a lot of money comes from our federal government and so a number of different federal agencies and some internal money from the state of Michigan through the university itself a number of graduate students over the years who have worked out on the details in this project some people have collaborated with Tom wheeler and Darren Mason I have no idea why those fonts changed but this this may be what happens when we change my talk on to your computer there may be some pride some surprises as we go some ups sorry some stuff will talk about the end with Kevin spear and pros pros at Case Western Reserve and then at the very end if I have a little bit of time I want to cover some stuff for looking at sort of mesoscale strain measurements with the new technique that we’ve recently developed so I guess the question comes about you know why would we want to use SCM for examining micro scale plastic def versus things like tem well first of all how many people here have ever made a TEM sample okay how many people have ever enjoyed making a TEM sample okay you know certainly very difficult time consuming I have a hard time getting students to do this you know we have to make a thin foil SCM really only requires us to prepare one surface we do have to do a good job on one surface but because there’s only one surface is certainly less prone to artifacts there’s more things that can go wrong when you make a TM foil that may affect your final results and we can also quite often do tests after we prepare the sample rather than prior to the sample which gives us a chance to see if the sample is good before we do for example mechanical tests it’s more difficult to do with the TM foil also sem maybe non-destructive it may allow us to test the same area through a sequence of steps and and see how the microsomal changes during the deformation processes rather TN typically what we have to do is de for material extract a sample and see what it looks like and then we’re done with that sample also TM you know has a very limited thin area and especially if the sample preparation is difficult it’s a very limited thin area SCM typically allows us to look at very large areas and this makes it much more useful to try to assess a broader scale nature of the material of interest and

also if we use bulk samples we can typically do in sit two experiments in an easier way I have done Institute tem and watch out dislocations moving in TM it’s very fascinating to do but for example the state of stress that you get when you do a tensile test on an S or on a TM foil is very difficult to characterize you really don’t know exactly what’s going on with a bulk sample in sem we have a very good idea of what the bulk status tresses and we can relate that more directly with what’s going on in the material so hopefully we’ll see some of those things as we go on but the other hand is you know sem is not going to solve all the problems we do have to use it in compliment to tem and sometimes we’re just stuck having to do to you so when I talk about these sort of techniques some of these will be familiar to you and I think some will be new electron backscatter diffraction patterns have become very common in the last 20 years we use them very extensively when I started doing them it would take 30 minutes to determine one orientation now with fully automated systems we can get 500 orientations a second so you can just get extreme amounts of data very quickly using it it’s very useful for orientation determination has significant but not perfect applications and phase identification well we focus on I’ll talk a lot about today for some different reasons our electronics handling patterns and they’ve been around for a very long time and sort of the subset of those are selected area camelina patterns they allow us to determine the orientation perhaps more accurately or in a different orientation or configuration than you do with electron backscatter diffraction they can also be used by phase I for phase identification but we’re not going to be doing that and then really the thing that is somewhat unique is this idea of electron channeling contrast imaging the idea that we can image fine scale defects such as dislocations and micro twins in sem something that a lot of people haven’t thought about doing in the past so very briefly can we see that well enough that’s pretty dark and I apologize I don’t know if it helps if we turn off the lights in the room will that hurt you but we can’t do that yes okay okay yeah there are actually very few dark lines on these slides so hopefully we’ll be okay whoops sorry so the idea with electron backscatter diffraction that we have a sample tilted to a fairly large angle we have an incident station or electron being we get inelastic scattering in the sample and those inelastically scattered electrons diffract form kikuchi lines that we detect on some sort of a screen and so the idea is you have sample again is tilted we focus the spot on a given area we collect these on a fluorescent screen which then is recorded with the ccd camera and we get patterns like this which we can then use a hustle Hough transform to determine the orientation and so we can do maps like this and actually this is not accurate now this is a few years old this is rates up to 30 orient all right 70 orientation per second this is more like about 500 orientations per second if everything is now working correctly but more we’re not going to focus on is this idea of electrons handling electronic channeling what we have the situation that simply as the electron beam orientation changes or if the orientation electron beam remains the same but the crystal orientation changes the backscattered intensity is going to change and so there’s a number of different ways that that can be used to give us information about our sample if we take a typical polished polycrystalline sample and we look at it in typical backscattered mode and if we don’t have other things going on like phase contrast and things like that grains with different orientations are going to give us different back scattered intensity so as a result we can see the grains quite easily you can even see little details like the sub grain boundary along here and there’s another one there that is indicative of this change in back scattered intensity or backscattered yield as a function of the crystal orientation if we take a single crystal of silicon so now the orientation remains the same but we look at that at very low magnification as the orientation of the beam changes as we image a sample and that’s actively what happens in an SEM at very low magnification we get a being raster where the beam orientation changes will get a series of lines that is indicative of the crystallography of the sample and each one of these bands here is going to represent a set of crystallographic planes in the sample so from this we could determine the orientation of the sample but in this format this is not very useful for us because we’re really limited to using this on very large

single crystals but if instead if we rock the beam back and forth through two angles and try to rock that on a point on our sample as we rock that beam of those two angles and we output the intensity we can actually then get a pattern that now reflects the crystallography of the sample and so the center of this pattern represents the unwrapped beam or the orientation of the microscope axis in this pattern now which we refer to as a selected area candling pattern this pattern here is going to represent what is the orientation of the sample relative to the optic axis of the microscope and that becomes very valuable for us before we’re done and we typically don’t use these necessarily for trying to existing general find orientations but we want to be able to use these in the way you would use a selected area diffraction pattern and TEM in order for us to set up specific imaging conditions and so that’s quite useful for us there are some problems with this in just in terms of the general sense of a scanning electron microscope what you can do with this first of all a lot of scms are incapable of doing this that used to be offered is a fairly standard option on sem s but it is much less common these days because ebsd has become so widespread use but first of all what happens if we don’t rock this beam on the surface of the sample let’s imagine out it’s rocking down here and that’s very common in a lot of microscopes well certainly then is e rock the beam is going to sweep across a sample and so if you cross things like grain boundaries or in two different orientations we’re going get channeling information from those different orientations and so this means that the resolution of the technique of space resolution is really quite limit so if we can adjust this rocking point and put it right on the sample we can certainly get information from a smaller area which becomes more useful for us but there’s still a problem here and if you look at both of these in these what we see is that actually these rocking beams don’t come to a point and the reason is is this beam rocks back and forth it’s subject to steer collaboration in the objective lens of the sem i kind of have it drawn backwards but generally what we’re saying is the beams that are farther from the optic axis or more strongly focused and a result of that you get a disk of lease confusion associated which really defines the minimum area that you could collect information from but if you actually just the strength of the objective lens as you rock the beam you can actually get that area down look smaller by using a dynamic correction and as a result we can get information from a smaller area on the sample and the smaller the area the more useful it is for us because it allows us to measure specific orientations and small grains or if grains have orientation gradients in them we can separate those out so if we kind of make a comparison of what an EV SD pattern looks like compared to a selected eric ameling pattern here’s an EBS be paired it actually collects a very large angular range probably about anywhere from 45 to 60 degrees depending on how you have the system set up set up so we have these bands that now correspond to these different crystallographic planes where a lot of these bands come together we get zone axes where that represent high symmetry orientations there are some limitations things like super lattice information is not contained in these bands it’s there but it’s really hard to see we compare that to secretary handling pattern here we actually collect a much smaller area now we can directly compare these two but now we’re collecting information from maybe five degrees and so that makes it more difficult to determine exact orientations because we don’t have all this broad information on the other hand we can see fine details for example these lines here are super lattice fans don’t show up in these other two bands this happens to be a 111 axis in titanium aluminum and now the center of this pattern again corresponds to the electron beam direction during normal scanning and that’s what makes them very useful for us so to kind of compare those a little bit more EBS vs space or resolution certainly less than 1 micron some people will say you know tens of nanometers it really all depends on how you’re defining your backscattered electron interaction volume and how that is a function of direction in your tilt to sample it has a fairly complex interaction volume the angular accuracy in general is not all that good first of all there’s scatter and how the Hough transform solves the orientations but on top of that what happens is this is a calibrated orientation it’s done based on putting a piece of silicon in the

microscope and calibrating that orientation and then hoping you get your sample oriented in the exact same manner and if the orientation angles are off a little bit as you put that in the should say the mounting angles that’s going to cause some problems there there are certainly some techniques out there for example angus wilkinson and oxfords come up with a relative orientation accuracy which is extremely accurate down to about point zero zero one degrees but that is really just relative to other orientations in the local area again fully automated though very fast fine structures not well resolved with select very handling pattern it’s about the best we can do is maybe about five microns so the space resolution is not nearly as good if we don’t have dynamic correction is going to be even worse about 100 microns which really starts limiting this approach but the angular accuracy relative to the optic axis is quite high it’s about a tenth of a degree maybe a little bit better only minimum automation which my students don’t like because it makes them work hard okay and we do get fine structure resolved which I’ll give an example here how that can be really very useful for us so if we kind of look at the sorts of things we can do I think can imagine now if we could see contrast from different orientations in the polycrystal we could certainly see contrast from different orientations associated with twins as we move across a twin when the orientation changes we’re going to change in that backscattered yield which is going to reveal the twin we can actually do that with pretty good resolution this is a secondary electron image have deformed ti al and this shows actually a twin boundary a kneeling twin boundary and then these fine lines are actually deformation twins and those are very very small you can see in some cases those are in the order of tens to maybe hundreds of nanometers across and so they’re very fine twins that we can resolve actually very easily in the sem if we set up our imaging conditions correctly if we want to see a dislocation imagine this as a dislocation here as we’re after the beam across the sample we can imagine now locally the orientation is going to change and as a result if we could imagine the camming pattern that we could get near that dislocation we can’t measure that obviously but but we can measure it we can imagine that the center of those channeling patterns will represent the local backscattered yield and that local backscatter deal will be given to us by what is the intensity on that pattern at that given orientation and because the patterns just as we move across the sample we’re get change as an intensity as the beam moves across that dislocation line and as a result again here’s a secondary electron and makes a very flat polished surface but when we imaged it under the right candling conditions if we set up the orientation properly we can actually see the dislocations in the substructure actually fairly clearly there’s actually a deformation twin you can see it sort of Peters out here and so these dislocations right through here are actually the dislocations at the leading edge of that twin so we can resolve them reasonably well maybe not as good as TM but not perfect now this is kind of noisy and make sure but i want to show you that you know funny you know people look at these things to say yeah well the dislocations look a little bit different and they did doing TM well you might expect them to because we have a bulk sample and we’re only really looking at the dislocations that are close to the surface in fact theoretical work was done in Peter hearses group many many years ago on this and so it’s basically if we take the trace across a dislocation line as you go deeper into the sample that intensity p the background is going to decrease and as a result we can imagine that dislocations are going to be sharp or the exit the surface and as i go deeper into the surface they’re going to become noisier and eventually disappear and at first people say well that’s sort of inconvenient but the reality is that’s a really good thing because if we’re seeing all the dislocations through the depth of the bulk sample we wouldn’t be able to see anything they would completely overlap and make it difficult to see so this is a useful thing for us but it also means that it’s going to be a little bit more difficult for us to interpret the dislocations relative to how we’re used to seeing them in tem well how do we do this well we it’s a fairly straightforward situation we use a field-emission gun sem we use a typical backscattered electron detector pole piece mounted on some of our microscopes we’ve gone to retractable detectors because we want to be able to get out of the way it at times what we do and i’ll discuss a little bit more but we generally use the specimen pretty much perpendicular to the incoming electron beam we can tilt it about plus

or minus 30 degrees and still be able to use this fairly usefully but this is a different configuration than a couple of other people use the field emission gun uses a thermal gun gives us a pretty good beam current that’s been a big advantage to us we have a beam divergence of about eight Miller radians we work at typical sem kind of working distances and that makes our life really nice because what we’re able to do is use various deformation stages things like that that would be much more difficult to do if we are in a high tilt configuration and quite frankly impossible to do if we are doing this in a TEM this typical back see our detector with these working distances gives us a fairly large detector collection angle which is very useful for us because what it does is gives us a lot more electrons on the detector which really makes life easy and one of the critical things for us is generally you’re going to need to be able to collect these channeling patterns in order to set up the proper imaging conditions just going through a number the parameters in no particular order one of the reasons you need to feel in the mission gun is you need a high brightness but you need to do that in combination with a small beam convergence in a typical thermionic either tungsten or lab six-gun really can’t do that for you there’s some theoretical work by David joy about 20 years ago that shows some dislocation profiles as a function of em divergent so you can see the peak the background decreases as a beam divergence increases and if you get much below this you’re going to get a real flat and you really can’t image your your samples well at all the other thing worth noting on here if you look at these sort of peak to background positions here the widths of the dislocation lines are similar to that that you’d see in brightfield tem imaging that’s about as good as we’re going to do we can’t set up anything like a week beam condition like you get in a typical tem because here we’re really limited by the probe size in terms of our resolution rather than some of the other things associated with week being microscopy how do we set this up well in many ways this is the same thing as you would do in a TEM and TEM what you do is you would go to the fraction mode you would tilt your sample to a to beam condition you’d adjust your deviation parameter by moving your coochie lines relative to the positions of your diffracted beams and then you switch back to imaging mode well what we do here is we go to the fraction mode we tilt 2a tubing condition we increase the magnification of that pattern and we put the center of our pattern right on the edge of one of these channeling bands and the reason we do that is this is where we get the most dramatic change in contrast between the dislocations of bending the planes in the background and then what we do is we just switch back to typical imaging mode so here we’re rocking the beam and then here we’re just rat stirring the beam back and forth across the sample so many ways we can envision it a lot like what we do in diffraction contrast tem one thing that’s a little bit different and those of you who are familiar with TM you realize that dynamical theory in particular will actually kinematical and dynamical theory both tell us that we want to use a deviation from the exact bragg condition and dynamical theory says we have to do that with a positive deviation from the exact bragg condition it turns out with this we actually want to set up our channeling condition exactly at the Bragg condition this shows again a number of theoretical profiles and again this was all thought about really 40 years ago by hearses group and they sort of gave it up because they didn’t really have good microscopes for doing it and so you can see that if our dimensionless deviation parameter is equal to 0 this shows that maximum peak to peak contrasts associated with this and so what I’ve shown here now are some experimental images that we’ve done this is a secular channeling pattern that effectively shows the locations of the non tilted beam which gives us really what the deviation parameter relative to this channeling band we can see when s is less than zero the contrast is very poor when s is equal to 0 we get strong contrast when this is greater than zero we maybe don’t get a strong contrast but we get a little bit sharper images and so generally we want to be in these kind of imaging conditions in order for us to be able to get the best image best contrast best resolution of our dislocations a couple of other things if you’re familiar with tem you can imagine setting up diffracting conditions with a minus g condition or a plus G condition the same sort of situation occurs we get a reversal in contrast we can see for example we go from here dark to bright here we go right to dark which really

make sense because we’re saying one side that this location is giving us a dominant diffraction rather than the other side of the dislocation we can also do things Qaddafi contrast analysis it’s not trivial because tilting from one bright condition to another Brad condition is time consuming and not as easy as it is in tem got a problem o light okay yeah light light would be fine here okay and again you know here’s an area where this is second your electron image either dislocations and silicon here and the critical thing is you can see there are changes in contrast with these different dislocations as we tilt two different handling conditions these are the only tem images i’m going to show you but but there’s a reason for this because i want to demonstrate the gdp cross you criteria what these are our prison attic loops in iron aluminum the burgers vector is 0 0 1 which is coming out of the board and so when we look at these dislocations here a and B those both have the same burgers vector the critical thing is depending on your D vector when your G vector is parallel with the dislocation line those dislocations go out of contrast and we can see that both here with the 0 to 0 and here with the 200 reflections when we go over and we do the same thing with the candling conditions we can see that in this image we get segments that are both light or dark that are both in contrast for these 110 type reflections but when we go to these 200 reflections we see only one direction is visible so again the same criterias hold and so that means that if we’re very careful in this very tedious work i’ll be honest my students don’t like doing it but we can determine things like burgers vectors and we can certainly determine things like lying directions through trace analysis rather than having to make a tem foil to do those things and again sometimes that’s a valuable thing to do i want to take a second to talk about how we do our sort of approach versus other people a lot of people actually do this at a fairly high tilt configuration and the reason is is your backscattered yield increases as your tilt increases and so typically ebsd has done at large angles and that’s the reason for doing them at large angles that and also you can fit a camera on the side of the chamber you can’t put a camera right on top of the pole piece of your microscope and so that certainly is an advantage because you’ll get more signal at higher tilts but there’s some disadvantages to that first of all if you’re at high tilts what tends to happen is your image just squished together in one direction and you have to focus differently as you r Astor across your sample you can certainly do some corrections in a microscope to take care of some of those things but it isn’t necessarily an easy thing to do this is the same area shown here and this is a lot easier to understand what’s going on here than it is here but the thing that we have found more difficult is when we are collecting these selected erics handling patterns if we do it at a high tilt what happens is the intensity is varying as we’re rocking the beam back and forth because that trajectory is changing so your selected area channeling patterns will come out something like this but if you do that same sort of em Rocking through say 10 degrees here we can see that we get much more even intensity and this just makes life easier the third thing is if we have our sample tilted way over it becomes much more difficult for us for example to use our Institute deformation stages they start knocking up against the pole piece of our microscope and things like that so just working at a low tilt makes life a lot easier in a microscope and so that’s why we prefer to do it that way and so given that it’s really quite this technique electronic sailing contrast imaging is really very well suited for particularly in sit two experiments I want I can go through is I’ll show you some things from both in sit to an exit to it experiments well the sort of thing that we can do is we can do both into or exit to four point bending the advantage here with four point bending is we can look at the tensile surface of the sample were the maximum stresses and that really helps us work our way through some things we can also look at the sides of four point been specimens institue and i’ll show you some examples is and this is quite useful for us to get a general feel for how a crack is propagating and finally what we can do is we can actually do some fairly high cycle fatigue and bending to which again allows us to for example at the tensile surface in a fatigue sample we have a stage that will go at about 500 cycles per minute on our microscope obviously we have to stop it to actually image but we can do some fatigue testing and then we can look at we do some more fatigue testing and look at it which is a useful for thing for us what I’d like to do is in August go through a few examples of sorts of

things that we can do with this the first thing I want to show you is how we can look at crack tips and this happens to being Nick aluminium which 20 years ago as a very exciting things for people look at we’ve been so looking at this a little bit more and more over the years we’ve done is we deform these with four point bending and we in this particular case I won’t go into a lot of details about this but i want to show you that really the capabilities of the technique we can determine how alloy purity and thermal history are affecting the crack-tip behavior and the crack propagation the reason we’re looking at this is a commercial purity Nick aluminum is susceptible to strain aging it embrittlement it becomes a sexually a zero fracture toughness material if it is a cool the wrong way high purity neck alone is much less susceptible these variations in thermal history and so what we’re doing is we’re taking a sample like this a nickel loom is still it is very brittle so we actually bind it to an aluminum substrate really just to increase the toughness a little bit we take the sample and we put it in our deformation stage here and you can see it right there may be okay and all we’re doing then is we’re deforming this and we’re looking at how the crack propagates out of the snacks here and so this shows you the knock you can see how the crack deformed here and actually branched here here’s the secondary electron image of that crack tip but now when we actually look at the crack tip with channeling conditions we can see the individual dislocations out in front of the crack tip so we can measure the extent of the plastic zone certainly as we get closer the crack tip the number of dislocations increases dramatically that is partly the reason we get all this blooming here also what happens is as the beam comes very closely eggs of samples some more backscattered electrons leaked out the side which gives us increase intensity but the critical thing is really what’s occurring out here in terms of being able to directly measure that plastic deformation filled out in front of a crack the other thing knows because we have a bulk sample instead of just looking at the crack what we can do is we can look anywhere along that previous crack and so we get a feel for how big the crack tip was as it propagated here we can see in the slowpoke commercially pure alley there’s just a couple of dislocations along there and you can see that really it’s a very flat fracture in the same high purity material slow cooled we can see it’s a much more torturous path I think there’s an indication of little bits of plasticity right along there well we can also see this location slip bands along here and we can see that those on this side go out to about here which gives us a feel for how big that plastic zone was as the crack was dynamically propagating now there’s maybe some relaxation after the Lotus is released everything but what we’ve seen in general is the crack tips are actually say the plastic zone ends up being much larger when the crack actually arrest so it runs out there and when it slows down we seem to get a little bit larger plastic field the other thing that’s interesting to note is how the dislocations are significantly different on either side of the crack and we have them characterize these fully but certainly these dislocations are more or less end on so they look like points or these dislocations here are lying in the plane of the foil and sort not foil plane of the sample and certainly aren’t lined up in these same sorts of slip bands along here and so I think we have a tendency when we think about practice to think about the dislocations being pushed out in a set metrical way and from the crack tip but the reality is it doesn’t need to be the case and if you really think about it it probably isn’t expected to be the case and so we can certainly see things like that okay I’m going to talk a little bit more detail about ti al this Agana alec and we’re spend quite a bit of time talking about deformation transfer this alloy and what we’re trying to do is really understand the processes of what happens at a grain boundary in terms of if we have a dislocation pile up or twin piling up at a grain boundary how does that transfer deformation into the next grain and other some grain boundaries that happens easily are there other grain boundaries where that is more difficult and particularly are there some grain boundaries that are susceptible to damaging through this process and so in a broad sense we’re trying to understand how texture and boundary miss orientations affect the ductility and fracture process and we’re trying to experimentally measure how dislocation between generation about occurs in these polycrystalline materials and then quantitatively characterize the impact of this boundary misorientation on this deformation transfer so to give you a feel and I don’t want to go into a lot of detail there are some people here who do know about ti al okay others of you may not be that familiar with it it’s an FCC

derivative structure but as a result of how the planes order we have super dislocations I’m sorry ordinary dislocations super dislocations and deformation twins and the deformation twins are kind of unique is that that they occur on the typical twinning plane and FCC the one on one planes but there’s only one twinning direction the other 220 directions won’t be active because they would be super lattice intrinsic stacking faults with much larger burgers vectors and so we get all these things going but one of the tricky things in here one thing that ebsd struggles with is because the sea on a ratio is very close here ebsd has a hard time distinguishing between the sea direction and the a direction but this leads to a large anisotropy in the plastic behavior so we really need to know the difference between those two and we can do that by identifying where these super lattice bands are which I’ve shown on this quadrant of a sara graphic projection so instead of having to deal with a simple stereographic triangle which you would have to an FCC we really have to worry about a much larger stereographic a portion of the stereographic projection give us all the information they’re so the first thing that we did is we said okay let’s try to map material like this with ebsd and what we find is we get a lot of spec Faline going on there and the reason is is one time it decides one direction is the a direction and another time besides another direction is the a direction is a hard time sorting that out and so what we found is if we use selected area channeling patterns we can now go in there and in particular if we focus on this zone here we can see in one of the three to 20 bands we can see the superlattice bands those superlattice bands then allow us to determine exactly where our orientation is within this stereographic quadrant using that then we can take this pattern and run a little algorithm to correct these patterns to give us the true orientations but then what we can do is not only do we know the true orientation of a particular grain what we can do is we actually then can go back and see the large crystal rotations that are associated with the plastic deformation there so this is something that evsc really doesn’t allow you to do we have to supplement the ebsd with the selected area channeling oh okay that’s what my students do for me I apologize for that and I didn’t know if those if I’d done the PDF that wouldn’t have happened on me okay and so then what we want to do is really look at the same a light but we’re going to look at the deformation behavior a little bit by looking at first of all this loading configuration which gives us in into what’s going on and then looking at this loading condition which is going to give us some statistical basis to make some conclusions and so the first thing I want to do is you know look at some of these boundaries and see what’s going on again these are SCM images and here we can see twins piling up into this boundary here and we can see actually large numbers of dislocations being transferred into the next screen associated with these twin pileups we can also see there certainly background dislocations here and here but we can see there’s a much larger really too large to resolve numbers of dislocations associated with that strain transfer there if we look at a crack prop getting out of one of these noxious here we can see some interesting things as the crack propagates along in general what happens is it propagates by cleavage but every now and then we get some intergranular failure and not only that we see places like this where the crack propagator here and suddenly jumped onto a different plane we’ve seen that a number of different places where we go from cleavage to intergranular and sometimes intergranular is reno creating on a different plane or I’m sorry sorry ahead of the crack okay so if I look at that crack tip in some more detail we can see there’s actually a lot of deformation going on here we can see that there are lots of deformation twins if we look closely here we see kind of a rougher what I call model background which suggests we have lots of dislocations there and they’re really difficult to resolve at this point because we know the true orientations of all these grains we can now do trace analysis to identify the specific planes and with those specific planes we know that there’s a specific twin direction so we know all the crystallographic details of these deformation twins you know exactly what twins they are rather than you’re saying we have a twin on a given plane and that’s going to become useful for us here in just a second what I want to do is concentrate on what’s going on up in this little zone up here which you’ll see in the next slide so this zone is right here what we see is we have a bunch of twins here and right at the tip of some of these twins we have some very small micro cracks that are nucleating and a first bus we looked

at it said would well the twins are running into the boundary and that’s creating a crack but in fact what’s happening you which i think is maybe even more interesting these twins are actually nucleating here and eat and the terminating you see some here terminate as they move out from that boundary so they nucleate here but when these twins nucleate they leave cracks behind this is kind of interesting and as we deform it a little bit more we find that actually that boundary right there opens up into a larger crack so this is a mechanism or what we have is deformation is blunting the cleavage crack but that deformation is also renew creating cracks further ahead okay so you know we tend to think of deformation is inhibiting fracture but deformation may also nucleate fracture okay and if we do the crystallographic analysis on this again we know the exact nature of those twins so you know twinning is uni-directional so we know we’ll have a specific shear associated with the twin we find that these micro cracks are located asymmetrically in relation to the twins and that asymmetric nature is completely consistent with the shear that we expect from those twins as they open up and so we’re very confident than I the micro cracks form as a direct consequence of the twins nucleating at that boundary there so we’ve done now is instead of looking at that side of a sample four point bending all we’ve done is we’ve gone and looked at the tensile surface of these four-point been specimens and what we find is there’s actually quite a few places in here where we get these sorts of cracks nucleating at these boundaries and so this is a little bit of a complex microstructure it’s not just pure gamma but we found that the majority of 20 boundary we saw crack the majority 15 min of those were at gamma gamma boundaries three were at alpha 2 gamma boundaries one of them was in between two alpha 2 particles and one was between a couple of lamellar colonies and so we’re going to focus on these right here where where do these boundaries crack and really the question comes up is why do they specific boundaries crack and other ones don’t you know we have twins interaction with the grain boundaries either good grain boundaries and bad grain boundaries and if we can understand what are the good grain boundaries and the bad but grain boundaries maybe we can do something and processing to eliminate the bad grain boundaries so I got in here are just some more examples of these we can see these little tiny micro cracks here the micro cracks are there and at first we looked at this in terms of a sub compatibility factor those developed by Maria Morris a number of years going to titanium aluminum which really just says how well aligned are the deformation systems what is the angle between the slip plane normals and what is the angle between the directions and it would make sense that if those are well aligned slip transfer might be easy and you won’t get deformation and so this factor here can range between one which says is fully compatible and minus one which would be fully compatible for dislocations but it would be anti compatible for twinning because twins are you need directional and you can’t share them in either direction okay a value of zero basically says that the deformation systems are 90 degrees to each other okay and so if we look at this in detail we find that and don’t have this on the next slide down okay we find that in general that if we have a high M Prime we tend to see that’s where we get evidence of transfer across boundaries maybe see if twin coming to a boundary and the dislocation pile up moving across or twin coming in and a twin propagating across and so the factor works really quite well in terms of deformation transfer question is if you get deformation transferred do you get fracture or do you not get fractured a little bit different question and so we looked at that we said okay what is this mid factor for a particular deformation twinning system and what is this M Prime and this shows it for twin twin compatibility but we’ve look at compatibility for different types of systems and we plotted the crack boundaries versus uncracked boundaries and you can see that there is not a strong correlation okay it says that this factor says that even if you have good deformation compatibility that does not predict whether or not you’ll get fracture at the boundaries so what we do is we came up with a different factor based on work by Gibson forward who sits at the twinning share in this material tends to be accommodated by dislocation motion and we put together this this factor here that takes into account a lot of things first of all the Smith factor of a given twin system in other words how likely is it for that twin system to be active how is the burgers

vector associated with that related to the tensile direction because if these are parallel with each other it means that in general you’re going to have more of an opening force on that boundary and then how well can the twin share be associate or it can be a related to ordinary dislocations here in the material in other words how well do these accommodate each other so this kind of shows a model of this where if we have cracking we’re going to have this situation if we don’t have cracking we’re going to have this situation taking into account these various factors so when we did that we ended up looking at a couple of populations a population of crack boundaries and a population of uncracked boundaries and I was going to kick out of this because now I’ve done tem for 30 years and nobody ever does statistical analysis using tem and why is that because how many people have ever made a tem foil ok but this allows us to do enough observations for us to get some relatively statistically relevant types of data and so we’ve done here is we’ve used this fact room we’ve calculated it for the crack boundaries and the uncracked boundaries and we looked at those populations and did a student t-test is familiar with the key tests that compares two populations and the question that it really asked is it a it is it actually one population or as a two populations and so based on this factor we do this t-test and we get a probability that these are actually the same population of 0 point 0 0 5 which suggests that these are actually distinct populations it doesn’t mean we can necessarily predict what boundaries are going to be cracking but we can also say if we take a collection of boundaries and we look at what their factors are is it more likely one falls in one population or more likely it falls in the other population which is getting us towards predictive ities the most interesting thing though is if we look at the individual factors associated with this we do the same sorts of tests we see that the probability doesn’t break out the same way so really what it says is that the boundary cracking we’re seeing here is a function of multiple variables that are involved and I guess it shouldn’t be a surprise anyone but it also means that whenever you go to material you’ve really got to look at a lot of the details and the crystallography and how things interact to really understand why some boundaries may fracturing why some boundaries may not fracture it’s not a simple process okay so I want to show up just a couple of other applications here a little bit of work we’ve done in terms of device materials five minutes okay five minutes I’m almost there okay dislocation densities tend to be lower in device materials which actually makes this technique maybe even a little bit easier to use this is an MOCVD gallium nitride these are threading dislocations again here’s a secondary or China is a extremely smooth surface but we get lots of threading dislocations coming through and so this gives you the ability really the map where the dislocations are quite easily but what we want to do is see if we could apply this a little bit more broadly and this is a stuff with peruse and Kevin’s feared a case western where they brought me this ship and they said well they’ve done some thermal imaging on this and they see these lines and they think they’re dislocation they said can we see if these are really dislocations I said well we can see if we have lines that image like dislocations and they look that way and a first blush we thought this would be quite a challenge because it’s a little bit hard to get an orientation from something like this or so we thought but the reality is is that this is a hetero epitaxial situation and this is an exactly the same orientation is this and this is an exact the same orientation as this and so it makes it actually pretty easy to get a channeling pattern we certainly get an interference from the various structures on there but we can get the pattern really without any difficulty and looking at that this is you know what they had looked at with their optical emission microscopy this is now in one of these areas here where we see exactly the same sorts of structures obviously much better resolution we did some tilting experiments and we were able to get some of the dislocations go out of contrast which helps confirm to us that these really are dislocations we didn’t do a full analysis but it certainly appeared that these are 16 to 1 1 type partial stacking faults that are lying in the plane of these structures and so we can look at it pretty easily okay so just as a summary and then I don’t know if I’ll do the last thing or not some folks you might be interested in okay sample preparation is certainly a lot easier for ecchi than it is for tem that should be a thin foil not a thing full only had that up there for a couple years okay it only requires us to prepare one surface

it can be not destructive I can actually look at it and see if I have dislocations then through the deformation and see how the dislocations actually appear and so that makes it quite useful for us we can look at very large areas you know if the FIB you can do some very nice site selectively TM foil preparation these days but you still have to take samples from a lot of different places here I can look all over the place without actually damaging the sample so there’s some advantages there so this makes it very conducive to in sit two experiments we have bulk samples more representative in terms of their state of stress we have a large s CM chamber so it gives us a lot of options we can actually now do this at we have an image dislocations yet but we can certainly do institue experiments up to 900 degrees C with straining so we can do lots of different things there certainly it’s going to be less prone to sample preparation artifacts ooh maybe won’t eliminate them we select a pair of surface we still have issues with image forces on dislocations but we can track those a little bit better and we can look at very specific areas of interest where tem that historically has been more difficult to do with a fib you can perhaps do that a little bit easier now but but this gives us some real advantages so just kind of in summary this really summarizes the other things that I talked about a couple minutes okay so what I’m gonna do is I’m just going to jump ahead is something that I know maybe interest civil engineers a little bit more because technically this is a school civil engineering isn’t it okay and so I got working with Ian Patterson when he was at Michigan State and Rachel Tomlinson she was on sabbatical from Sheffield and they’re doing fatigue tests in compact tension specimen and they have crack tips and I said oh I can look at crack tips well it turned out that their crack tips had a lot of dislocations okay and we couldn’t really image them specifically but what we end up doing is we said okay what happens if we look at this whole area out in front of a crack tip and so this is very large I mean compared to what we would typically look at with the things I do this is a huge area but they’re very interested in what is the size of the plastic field here and that is something that has historically has been very difficult to do they’re using thermal imaging during fatigue to actually measure that they wanted to know if they were right well maybe right okay so what we do is we collect the whole bunch of backscattered images there and that’s really simple to do just go in there have an undergraduate snap image after image but what happens is the images near the crack-tip look different than the images far away because near the crack-tip there’s more dislocations there’s more crystal rotations and the channeling contrast is going to be different in those areas so as a consequence again if you move far away you can see the rain structure in these backscattered images fairly easily but close to the crack tip because of the crystal rotations we get more channeling bands and things like that it becomes harder to see those and what we said is if there’s a change in this contrast like this maybe what we can do is do a Fourier transform on that and when we do that it was very whoops sorry about that it was very clear that the Fourier transform changed because the frequencies in the images were changing again these are fairly low magnification it makes as that’s 100 microns there and if we take a trace across those we can actually turn around and measure the full width at half we do a bit of cleaning up to do that and which now gives us something that says okay the full width at half max is going to change as we deform the material so can we quantify that or quantify that we didn’t insert to straining test of the exact same material and we measured the full width at half maximum ages over these different strains so we developed really what’s just a calibration curve here’s a foe with it f max for various engineering strains and by doing that then we were able to take this crack tip area and again here’s one millimeter and actually map out what is the strain in front of the crack tip now you know microscopy us and everyone keeps working smaller and smaller okay but this is macroscopic this is actually fairly easy to do and in like i said i think maybe civil engineers appreciate this when i saw this is a mechanical engineers they get really excited because this is a scale there that they’re interested in and was kind of amazing um as far as we could tell no one had ever taken this kind of approach before we recently published it in proceedings of the royal society so they must have liked it and we were able to go back and then make a comparison with the von Mises the dugdale knee erwin plastic zones and we can see that you know it’s actually pretty consistent and

so i thought i’d throw this into something that may be a few folks around here might be interested in it’s really exactly the same approaches as we’re taking on the microscopic level really add a much more mesoscopic level so i think that’s your slide ok thank you very much I’m f it today is a holiday for him I have to take him to class so you have I have two and a half hours of lecture today yeah we work a lot hmm every day I know siesta right we have time to get for two very easy question I hope I’m rocking the beam yeah okay um it depends it used to be a lot of microscopes offered it standard okay because you can use your scam coils to rock the beam if you change the strength of some your other lenses but typically that rocks at a large working distances the Joel a 40 series would rocket 40 millimeters and then there’s other problems with that ok so our microscope we have two microscopes they can do this actually have a dedicated set of rocking coils which makes life easier and so a lot of people ask me could I do this instead of with circular candling can I do with ebsd and my answer is maybe ok you know Stefan’s F here at max planck and distal Dorf ok he’s an ebsd guru and he has a program that allows him to get an orientation with ebsd and then predict what tilt see need to get to what various handling conditions ok the problem is if the ebsd patterns only actor to about one degree that is going to cause some problems but they what they do is they get close and then the exhausted tilt slightly until they get nice contrast and so I think it works fairly well the other thing that I’ve thought about and we may do this because we want some things in our dual beam fib which doesn’t have this okay is if you co mounted a piece of silicon to do your orientation calibration and then got a candling pattern at low magnification which you can do in any microscope on that silicon you could calibrate the relative positions between non tilted and tilted which would help you with that accuracy yeah better than tem okay what we’re doing and I talked about this a little bit yesterday we’re doing some 3d x-ray at argonne okay which allows us to get a three-dimensional orientations okay which helps us but we also you know the are lousy patterns that we get allow a pattern streak from dislocations and if you assume the dislocations are edge dislocations you can determine what the streak directions should be it’s not as clear if they’re not edge dislocations the person does this Rose a bear a bath she says oh don’t worry about screw dislocation I worry about screw dislocations okay but at least in every case that we’ve looked and we’ve looked at about a half dozen from the dislocations that we’ve characterized based on trace analysis the streaking we get and three the x-ray is consistent with that so is it perfect no other surface relaxations yes it’s moving us forward okay what we’re doing and I talked about the Sun yesterday it is we’re taking we’re deforming some various microstructural packages we map them with ebsd we do a little bit of 3d and then we’re turning around we’re simulating those with the guys at Mach splunk using crystal plasticity and you know from a general sense they match up pretty well where we get in trouble

there is the details at grain boundaries because the cpf m fe m doesn’t handle grain boundaries very well but it may also be that we’ve got problems with our experiments because I don’t know what the inclination of a boundary is you know that was that geometric factor i gave is only based on this orientation doesn’t say anything about the boundary orientation so you know these are hard problems and maybe we’re making a little bit of progress but we’ve got a long ways to go which is good because I want to keep working for the next few years yeah okay off to class you