Severe Plastic Deformation - Science topic

Identifying GP zones (Guinier-Preston zones) in AA7075 alloy using Transmission Electron Microscopy (TEM) images requires an understanding of their characteristic features and the specific microstructural changes they undergo during aging. AA7075 is a high-strength aluminum alloy primarily used in aerospace and structural applications, and its microstructure is influenced by the precipitation of various phases (such as η' (MgZn₂) and η (MgZn₂)).What are GP Zones? GP zones are precipitates that form during the initial stages of aging of Al-Zn-Mg alloys, like AA7075. They consist of fine, coherent clusters of solute atoms (mainly Zn and Mg) that are in a precursor state to the more stable η' phase. These zones are often small, coherent (meaning they have a similar crystal structure to the matrix), and can be difficult to identify at early aging stages. Steps to Identify GP Zones Using TEM: 1. Sample Preparation: Thin Foil Preparation: For TEM, you'll need to prepare very thin foils (around 100–200 nm thick) of the AA7075 alloy. This is usually done using a focused ion beam (FIB) or mechanical polishing followed by ion milling to create the electron-transparent area. 2. TEM Imaging Conditions: High-resolution TEM (HRTEM): To observe the fine details of the microstructure, use HRTEM, which will help you visualize the lattice structure and any precipitates within the matrix. Electron diffraction: TEM-based electron diffraction can help identify the crystallographic nature of the precipitates and matrix, which can be crucial for identifying GP zones. Bright-field imaging: Useful for general structural observation. It can show the presence of precipitates as dark spots in the matrix due to differences in electron scattering. Dark-field imaging: When combined with specific diffraction spots, dark-field imaging allows for enhanced contrast of fine precipitates. 3. Characteristic Features of GP Zones in TEM: Size and Shape: GP zones are typically small, ranging from 2–10 nm in diameter. They may appear as fine, needle-like or spherical structures within the matrix. Coherent Nature: The GP zones are coherent with the matrix, meaning they maintain a continuous crystal lattice with the aluminum matrix. In HRTEM, they often appear as periodic dark and light contrast bands or small, circular regions. This contrast is due to their atomic misfit with the aluminum matrix, but the misfit is very small because the zones are coherent. Weak Contrast in Bright-field Imaging: GP zones are typically too small to be seen directly in bright-field TEM images, as their size is below the resolution limit of conventional imaging. However, they might appear as faint contrasts in the matrix due to small differences in electron scattering. Diffraction Contrast: In diffraction patterns, you may observe weak spots or streaks corresponding to the superlattice reflections associated with GP zones. These reflections typically appear at low angle, especially in the [001] or [011] zone axes of the matrix. Precipitate Distribution: GP zones can form in clusters, distributed in a semi-ordered manner within the matrix. In some cases, these clusters may be aligned along specific crystallographic directions. 4. Aging Time and Temperature Considerations: GP zones are typically observed at early stages of aging (during the first few hours to days of aging) before they evolve into larger precipitates like η' or η phases. The characteristic GP zone contrast is often most prominent in the early stages of aging, and the size and number of zones can increase as the material is further aged. In longer-aging conditions, GP zones may evolve into larger, η' precipitates, which are more readily observable under TEM. Thus, it's important to be aware of the aging condition of the sample and correlate it with the expected microstructure. 5. Other Techniques for Confirmation: Selected Area Electron Diffraction (SAED): This can be used to detect specific diffraction patterns characteristic of GP zones. The diffraction spots corresponding to the superlattice structure of the GP zones will provide a clear indication of their presence. Energy Dispersive X-ray Spectroscopy (EDS): EDS can be employed to analyze the chemical composition of the precipitates. GP zones, being rich in Mg and Zn, will show up as regions with a distinct elemental signature compared to the aluminum matrix. Key Steps for Identification: Examine the TEM image at a high magnification (around 500,000x or more). Look for fine, coherent precipitates, especially if you're analyzing a sample that has undergone an early aging process. Look for faint contrast variations within the matrix. In some cases, GP zones will appear as regions of subtle contrast that differ from the matrix, but this can be difficult to distinguish from the matrix itself at low magnifications. Confirm with diffraction patterns: Use selected area electron diffraction (SAED) to confirm the crystallographic relationship of the precipitates with the aluminum matrix, looking for superlattice reflections associated with GP zones. Use dark-field imaging: If available, dark-field imaging using specific diffraction spots can enhance the contrast of GP zones, making them more visible. EDS analysis: Perform energy dispersive X-ray spectroscopy to detect the Zn and Mg concentrations in the suspected regions. GP zones will exhibit a higher concentration of these elements compared to the surrounding matrix. Conclusion: Identifying GP zones in AA7075 using TEM requires a combination of high-resolution imaging and diffraction techniques. While they may not always be immediately apparent in bright-field images due to their small size and coherent nature, the use of HRTEM, electron diffraction, and EDS analysis can provide conclusive evidence of their presence. Additionally, understanding the aging conditions and the specific evolution of these zones into more stable phases (such as η' or η) can help you interpret the TEM images more effectively.

Dear Bal Mukund Mishra

As you know that the high stacking fault energy is the energy required to induce a higher stacking fault, which is a sort of significant defect in the crystal lattice that leading to disrupt the stacking sequence of atomic planes. For that, performing EBSD is limited in this case. My best regards …

Given that the packing densities are dissimilar (0.74 for fcc compared to 0.68 for bcc) and the "atomic radius" remains essentially unchanged for a given element irrespective of which lattice structure it forms, I would argue that volume-preserving tetragonal shear cannot accomplish such a transformation.

Is the proposal still on ?

The application of this method in the machine-building industry is economically very profitable. Therefore, this method is preferred in the automotive industry.

Hello,

I believe that science is moving from continuum mechanics to particle mechanics. Here, depending on the size of the particles, I expect major breakthroughs across the whole range of particle sizes. Up to now, even particle mechanics in the micro to centimeter range is mostly modeled by continuum mechanics. Of course, with respect to the rapid development of DEM methods. This will allow the modelling of particles with respect to their sizes and the digitalization of the field. I wish you a lot of success in finding new properties of particles and their collectives.

Prof. Dr. Jiří Zegzulka, TU Ostrava

Dear Akash,

considering the process of FSW, a totally new micro structure ist created.

So, the number of high-angle grain boundaries will be different from the base material anyway.

In Conclusion, fully recrystallised grains cannot be detected looking at the amount of high angle grain boundaries but by the dislocation density inside the grains (see definitions of "re-crystallisation" in textbooks of Metal Physics).

Best regards - Torsten

Check to see if Poisson's Ratio values <0.5 will solve the issue.

I am sorry

I work on coatings that are very porous. I have the bulk 8YSZ materials and method with me so I can do few tests if time permits.

Magnesium alloys are corrosion resistant and etching is based on differential corrosion response shown by the grain boundaries. Freshly prepared etchants and also freshly polished specimens are used to obtain better results. Increase the etching time or re-polishing after etching and then etching can be performed.

I have only downfeed velocity as boundaty condition mentioned in the journal. And how to give both rotation and downfeed to the feedstock? Do I need to use any predefined field in abaqus?

Besides saving meshing effort and computing time, the 2D axisymmetric models show the complete solution, whereas the 3D solution is seen on the surface. Only after cutting the 3D solution can be seen on individual sections (again this is only a partial view). However, also the expected solution must be symmetric. Often a non-symmetric solution of a symmetric problem matters. Typical examples are instabilities (like buckling).

Kiarash Mashoufi, thanks, your answer helped me anyway and I found the meaning of Mr zhilyaev comment by myself!

Hello Reza Darvishi Kamachali and Kiarash Mashoufi

I'm new to the drawing process and I'm looking for the software Deform 3D to conduct some analysis in my research.

I would like to know where I can find the software to download for free?

Someone here has the installation file of Deform 3D to share with me?

Thanks in advance!

Hi Mr. Bahmani. I recently investigate the effect of grain size on the corrosion rate of low carbon steel, I hope it would be helpful!

It is possible but you should simulate the process that you want to predict. I recommend Deform 3D software for that purpose.

Hi,

The grain size may influence the biocompatibility, so the alloys co-exist better with human tissue. Also, strength is always important for the implants. SPD may give rise to better properties on those aspects.

You can follow this article for detailed restoration and recrystallization mechanism

Hello,

It is very difficult to distinquish between CDRX and GDRX using OIM data of final processed materials. This comes from the fact that GDRX can be categerozied as a type of CDRX mechanism. Both of them inherit the deformation texture components. However, if you can see the transition microstructure (between BM and final microstruture), you can do as follows. Calculate the point to point misoriention along the shear or deformation bands inside the transition area. If the misorientations are low from point to point (lower than 15 degree), it means that GDRX is governing the grain structure formation. ...

I have used this method in my papers. You may use them as fallows:

1)

Heidarzadeh A, Saeid T, Klemm V, Chabok A, Pei Y. Effect of stacking fault energy on the restoration mechanisms and mechanical properties of friction stir welded copper alloys. Materials & Design. 2019 Jan 15;162:185-97.

2)

Heidarzadeh A, Saeid T, Klemm V. Microstructure, texture, and mechanical properties of friction stir welded commercial brass alloy. Materials Characterization. 2016 Sep 1;119:84-91.

Aditya Prakash , DIC is not necessary; if the slip traces are well developed you will be able to see them in EBSD maps directly - eg. in IPF maps (Fig. 3 of the publication), and more visibly in local misorientation maps (LAM, KAM etc.) This is shown in Fig. 7 along with the superposition of the poles.

DIC is only used here to image and count a large number of slip traces in a shorter time than EBSD.

Dear Ali,

As you know, we say CDRX because there is not distinct steps of nucleation and growth. During CDRX, whole of the microstructural evolution is continuous. So, it can be occurred by transformation of LAGBs to HAGBs, fragmentation, GDRX, etc. You can consider the GDRX and fragmentation mechanism as some types of CDRX. The term "fragmentation" some times has been used instead of CDRX in literature.

Please note that in most cases it is very hard to distinguish the different restoration mechanisms. For example, it is well developed that CDRX does not inherit the deformation texture. However, I found in my research that the existence of deformation textures in the material can not always be the sign of CDRX.

I think the story of dynamic restoration mechanisms have not been completed, yet.

Hello Ali

read the artical in link

if you have 2 distinct peaks in the raman spectra especially at ~1342 and 1578  cm-1 region which could be correlated to the presence of the D and G bands respectively. The D band is related to the degree of disorderness while the G band corresponds to the presence of the graphitized carbon.

In most severely deformed ultra-high strength steels, grains are of anisotropic shape and the following numbers correspond to the smallest dimensions. To my knowledge, the smallest grains are less than 10 nm and are obtained in severely cold-drawn pearlitic wires, that are also the strongest steels known today (doi.org/10.1103/PhysRevLett.113.106104). By high-pressure torsion of a martensitic low carbon steel, we have recently obtained about 30 nm (doi.org/10.1002/adem.201800202 , a more detailed analysis of the as-deformed state is under preparation). For bainitic steels, see for example doi.org/10.4028/www.scientific.net/MSF.584-586.655 and references therein. In my opinion, the challenges are to realize deformation conditions, which allow severe grain refinement without fracture, on the one hand and to have tools, that are not deformed themselves during severe plastic deformation (especially if you combine it with solid solution strengthening with higher carbon contents in medium/high carbon steels), on the other hand.

The VPSC7 manual can be found at the following link:

Thank you all for your enlightening answers. For the same question, I found the following paper by Cotterill very interesting. You can have a look.

Thank you for sharing your wisdom with me.

Dr. Darren Paul Broom and Dr. Rk Fakher Alfahed

I appreciate your kindly helps. Also, I decide to discover a new high entropy alloy for this research.

regards,

Sina Ghaemi Khiavi

 

in magnesium alloys cross slip can occur in two conditions : first when you increase temperature the amount of SFE increases and it will be possible to cross slip in both basal and prismatic planes and it is worth mentioning that by increasing temperature non-basal slip systems would be activated and cross slip becomes easier. The second fact is that if you add rare earth element to magnesium alloys, these elements would decrease the ratio c/a and this ratio will become smaller and close to 1.623 which is the ideal ratio for the tetragonality factor in HCP systems, hence non-basal slip systems will be activated and cross slip would be possible.

Dear Shokouh Attarilar ,

For metals and metal alloys, elastic constants primarily depend on the interatomic forces. They are surprisingly independent on composition and the extent of the plastic deformation applied. However, in cases where textures can be expected, like Ti and its alloys, anisotropy should be invoked due to severe plastic deformation. Further, for mixtures (think of certain ceramics and/or products obtained by sintering), the rule of mixture may apply. For non-metallic materials like thermoplastic polymers, the Young's modulus depends on the distance of the polymer chains and the nature of their bonds.

The attached link provides a way to an introductionary discussion of the above topics.

Dear Sina, before u go into that calculation u need to improve ur data quality. As shown in the attached picture ur max intensity is less than 500 cts.

Ur sample is powder or bulk? what is the step scan and scan time?

It can be calculate by PM2K; the software that implement WPPM as mention by Eduard.

To do so u need high quality data and SRM scan to get the instrument broadening.

U cud contact Prof Matteo Leoni for the free software (Maybe the FPA already implemented in the software but this also need the details of inst. setting)

Yes, you can do it easily as by choosing time step and its sub-steps. As per your requirement you can choose the time interval by putting the time sub-step size. Choose you analysis as transient i.e. time dependent.

It is quite common that increasing strength results in reduced ductility because of a reduced strain hardening capability.  In your case, a high dislocation density and small grain sizes could be the key factors responsible for the mechanical behavior that you mentioned. There have been many publications on how to appropriately design ultrafine-grained structures for high strength and good ductility. Some of these papers are listed below:

1. Y. H. Zhao, et al., Simultaneously Elevating the Ductility and Strength of Ultrafine-Grained Pure Cu, Adv. Mater. 18 (2006) 2949-2953

2. Y. H. Zhao, et al., Tougher ultrafine-grain Cu via high-angle grain boundaries and low dislocation density, Appl. Phys. Lett. 92 (2008) 081903

3. Y. H. Zhao, et al., Determining the optimal stacking fault energy for achieving high ductility in ultrafine-grained Cu-Zn alloys, Mater. Sci. Eng. A 493 (2008) 123 – 129

4. Y. H. Zhao, et al., Simultaneously elevating the ductility and strength of nanostructured alloys, Adv. Mater. 18 (2006) 2280 -2283

5. Y. M. Wang and E. Ma, Three strategies to achieve uniform tensile deformation in a nanostructured metal, Acta Mater. 52 (2004) 1699-1709

I am not sure if there could be a straight comparison like this, or if the two can be isolated indeed. For the cross slip to take place, that is forcing the dislocations out of their original slip plane, the slip must be saturated at least locally. So the work hardening until that state is provided by the increased dislocation density and their pile-ups due to their slip. And if cross-slip does start it is actually softening mechanism as it provides a way for the slip to continue.

On the other hand, if we consider a hypothetical case where the slip of an individual dislocation (where there is no considerable dislocation density, i.e. the isolated dislocation is relatively far away from the other dislocations) and the cross slip of a pair of partials constituting a stacking fault, the stress required fro the cross-slip action is greater than that for the slip for the same system considered. 

Magnesium is in the form of very anizotropic hexagonal cryslal lattice. The dynamic precipitation influeces dislocation sleeps. Therefore precipitation  induces both i.e. final crystallographic texture and strengthnes.

More so that SPD structures are particularly saturated with defects and thus even faster recrystallize than after regular plastic deformation. Just try to heat your SPD sample to 150C for about ten minutes and observe what then happens with grains...

Plastic deformation occurs above the yield stress by definition. In a tensile test at a given temperature, the strain rate is imposed and the stress and deformation (strain) vary. There is first an elastic (reversible) deformation under the yield stress and then a plastic (irreversible) deformation above the yield stress up to rupture. You can indeed encounter microplastic behaviour with local plastic deformation occurring at a stress under the global yield stress (of a test sample, in a component, etc.). But this is because the local stress is higher than the local yield stress.

Creep deformation can occur under the yield stress at a given temperature. In a creep test, the stress (or load) is imposed and the strain rate and deformation (strain) vary. In metals, after an instantaneous strain, the creep strain will show 3 stages called primary, secondary and tertiary creep. Tertiary creep ends with rupture. There is a threshold temperature for the full creep curve to be obtained. This temperature is around 0.4 Tm (Tm = melting temperature) with both temperatures in K. Below this temperature, only primary creep is observed. Above this temperature the three stages are observed up to rupture.

Ferrite and austenite have very different behaviours during plastic deformation. Ferrite presents dislocation cells, while austenite presents slip bands. Strain hardening also occur in different ways in both phases, and those are probably the root causes for the shear bands you observed.

Latest version of Deform  V 11 software solves the issue.

There is a method for observing the motion of individual dislocations during plastic deformation using x-rays. Here is the reference:

Measuring texture changes does not constrain plasticity models sufficiently to unambiguously deduce active slip systems.

I would suggest to play with contact conditions, it looks like there is no (sufficient) pressure

See my articles inpersian

Hello, Rohit,

The grain size influences the areas of grain boundaries. However, there seems no direct relationship between SFE and grain size, because the dissociation of a full dislocation takes place within the grain.

The grain size does affect another deformation mechanism-twinning, though. One needs to take care that fcc and hcp structures may have different properties assciated with SFE and twinning.

We need to know that the plasticity depends much more on generalized stacking-fault energy (GSFE) rather than stacking-fault energy. The maximum point of the generalized stacking-fault energy, or say unstable stacking-fault energy(USFE), dominates the width of a extended dislocation according to a modified P-N model. Well, if the width is too large, i.e. the USFE is too small, the cross-slip will be harder to be performed due to the pinning of dislocations. Therefore, the plasticity lies on the competition between the USFE and cross-slip probability which may have an equilibrium point to maximize the plasticity.

Best,

Yifeng Wu

This is a quite complex question. The correlation of crystal directions and tensor properties is straightforward and definitely described in any book about physical crystallography or material science. In order to use it you only need to become familiar with tensor mathematics. Time and place are here much to small to explain this.

The problem related to EBSD is a bit more complex since even if you know the tensor property (which is not always the case for any phase) you only measure the local distribution of crystal orientations. Except of MTEX (free downloadable Matlab toolbox) I myself do not now any other software which is able to "translate" an orientation distribution from EBSD data into a prediction of the property anisotropy. However, this is the actual problem somebody has if he is using EBSD or a comparable technique. The interpretation from EBSD data displayed as orientation map, pole figure or whatever is only in very rare cases simple so that many maps shown somewhere look perhaps nice, but that's it already. In so far your question is touching a very sensitive point. The selection of PF (which lattice plane or lattice direction distribution should be used) is not easy as well since it depends on the symmetry and the specific tensor properties. Since we often have no idea how the texture affects or material property we simply display some kind of standard map and use "standard" pole figures. From my feeling as standard the directions of symmetry elements might be useful but there is still nobody who gave a statement regarding this problem. Following this approach, for cubic {111} {001} and {110} are automatically selected, but these pole figures do not have a specific meaning for other symmetries. For hexagonal axes {100} and {110} are even equivalent.  Nevertheless, all these pole figures are only the expression of a specific orientation distribution so that all pole figures contain in principle exactly the same information but they don't show it in the same manner.

With IPFs it is even worse since the used IPFs are actually reduced IPFs since an IPF is a pole figure displaying e.g. the standard reference directions (poles) of X, Y, and Z with respect to the locally discovered crystal orientation. The reduced IPF only displays the distribution of  one reference direction taking into account the symmetry in the corresponding Laue group. Using the highest symmetry a Z parallel to (123) and (213) are assumed to be identical and will be described by {123} and encoded by the same color. This is not really correct since Euler descriptions using rotations only are not able to transfer both directions into each other so that regarding the crystal orientations they are not equivalent. But this is still another topic.

Anyway, the use of IPFs makes mainly then sense when you assume a correlation between a single loading direction applied. Then you are interested in the impact of your loading direction, i.e. parallel to Z, on the development of your orientation distribution. However, in order to use this for calculations (which should always be the goal) you need the orientation distribution which is not the same as pole figures or IPFs. As I told in the beginning, it is a quite stony way which is certainly the reason that commercial EBSD software only rudimentarily supports you. There is still a lot of work to combine locally observed crystal orientations and material properties. Pole figure and IPF are simply tools which enable to decide "stronger" or "weaker" but often do not tell you anything regarding the actual problem behind. Typical example is W (elastically isotrop). You can have a strong texture, but regarding the elasticity you won't see any anisotropy. So we are concentrating on  fields which seems to deliver more direct correlations  like the distribution of boundaries or their misorientation distribution. If you investigate a single crystal you can try it by selecting the concerned lattice planes and directions but even then an interpretation is not straightforward. For polycrystalline materials there are simply to weak tools available. Definitely, they can be derived since by the crystal orientation data you have at least this information. You only need to combine it in a meaningful way with the crystal property.   

From the point view of NDT, If the grain changes could make differences in the electrical conductivity or magnetic permeability, you could test eddy current, but you will need prepare special specimens for the calibration.

Cryorolling is one of the emerging severe plastic deformation method aimed to obtain bulk-nanostructure grained microstructure. Rolling metals in cryogenic temperature effectively suppresses the dynamic recovery, which leads to a microstructure containing several dislocation sites. Upon thermal treatment like annealing, these dislcations anhilates and polygonises into sub-cells of nano-metric size. 

Prevention of recovery in as cryorolled condition leads to extreme strain hardening of metals like aluminium resulting in their high strength and low ductility. By careful control of annealing process, a high strength cum high ductility material can be produced. Ageing can also lead to development of material with such properties.

Follow the works of C.C Koch to understand the mechanical behaviour of nanostructured material. Follow works of S.K. Panigrahi for cryorolling of aluminium alloys

Dear Polkowski,

The following paper on the tensile properties of Ni samples fabricated by SPD might be useful.

thx

Critical stress to form twins is rather high. That is why a notable stage of work hardening should be passed first. Nothing else. 

Dear Mr Sachin,

You can do groove rolling at VSSC, a kind of SPD process for making thin sheet.

Regards

A.G. Rao

Hi Mehdi,

I do not think that you should use the element deletion feature in this case unless the material really fractures in the process. Instead You should use the ALE - adaptive meshing feature to control the FE-mesh distortion in the process. 

BR,

Timo

It seems you defined ansiotropic plastic behaviour (Hills's function) but did it incorrectly. I would recommend to switch to something simpler - start with ideally plastic, then plasticity with hardening and then switch to aniisotropy (if this is what you want) to see where exactly the problem lies.

Perhaps a look into the inp-file material card may also show whether you unintentionally clicked some options.

Dear Reza Abdi Behnagh

the occurrence of  twinning frequently have been reported during SPD of Mg alloys such as ECAP, HPT and ABE processes. this phenomena is related two complex state of deformation during spd process.  ....in the case of twinning , we have 3 kind of twin in Mg alloys including tensile twin, compression twin and double twin among them tensile twinning will active in early stage of deformation.....it has been reported that for example compression twinning occurred in final stage of deformation and is responsible for fracture in mg alloys.   

I know that AC2T in Austria has such equipment. We have done some hot hardness there up to 700 celsius.

Kumar and Bombac show you the answer: the cold forging is possible according final shape complexity and forging machine available. If you want you can explore with FEA simulation.

I second what Nilesh mentioned:

For soft material you can fabricate it with a CNC.

For steel it is better if you purchase it.

I worked with a 6.4 mm diameter tool and had hard time machining it due to the small feature sizes. Ended up buying ready tools.

Experimentally it could be a bit difficult.

But theoretically it is not that challenging. You need to monitor the so-called "Neutral Plane" movement during the process. This is the location in your deforming material that separates compression and extrusion phenomena. Please note that this location is dynamic throughout the process, so the share of compression and extrusion on the overall deformation is changing during the process.

For detailed information I can refer you to our paper:

Ultrasonic Impact Technology (UIT)  can be employed, as UIT also delivers stress relief and grain modification besides that compressive stress.

Dear Dr.  Basheer,  Your question is not well defined!   However, we can make the following statements:

The dislocation multiplication  ( increase in dislocation density) occurs during the plastic deformation  through the activation of  the various dislocation sources such as Frank-Read Sources,  large precipitate (inclusions) particles having sufficiently big sizes that can't be  cut through by the moving dislocations but rather go around and leaving dislocation loops.  There is relationship between the dislocation density and the plastic shear strain, which is given by   d= d_o+  a x Gamap^m, , where  d_o  is initial (grown in) dislocation density  [metals: about  10⁸ cm/cm³,  and ionic and covalent crystals 10³ cm/cm³ ]  a  is a constant ( f metals:  about equal to  10⁸  cm/cm3), m is about  =1.

The applied tensile stress) required to induced plastic flow  [ which corresponds to about  0.2%  elongation for  the most FCC metals and alloys )    is known as  the elastic limit  or  yield stress. In the  elastic region  dislocation density variations  are almost negligible, one can only talk about their rearrangements if the temperature is sufficiently high enough to allow dislocations  to  climb and cross slip to form polygonization or low angle grain boundaries etc. Best regards. 

maybe this helps:

page 20 (an introduction):

Residual stresses can be evaluated from displacements of atoms at a microscopic scale. The changes of the distances between atoms correspond to shifts of Bragg reflections, which can be measured accurately by X ray or neutron diffraction. Collimation allows to probe separately small volumes and in this way to obtain a complete map of all the displacements in material. Neutron scattering has the advantage of a large penetration, due to small absorption by most of the elements. Instruments are permanently dedicated to this kind of studies in many neutron facilities. The principle of the technique is relatively easy and does not need very high fluxes. Because of its interest in metallurgical industries, the instrument becomes an attractive technique to be developed around medium flux research reactors.

(J. Teixeira, M. Ceretti)

best wishes from Vienna

Andi

ps.:

I just realized the date (April 2013). Probably I came too late :-)

What is the question here?

It is straightforward to show that the slip systems you mention exist, simply by considering the number of possible permutations of the indices. For the {110}<1-11> systems, for example:

(110)[1-10] (101)[10-1] (011)[01-1]

(110)[-110] (101)[-101] (011)[0-11]

(1-10)[110] (10-1)[101] (01-1)[011]

(1-10)[-1-10] (10-1)[-10-1] (01-1)[0-1-1]

Note that six of these are simply the other six running in the reverse direction. Whether you consider there to be twelve {110}<1-11> systems or six is a matter of convention.

(Note also that while these systems may be physically distinct, only five of them can be truly independent. For proof of this statement, see Kelly & Knowles, 'Crystallography and Crystal Defects'.)

If you are asking whether all 48 systems can be activated in any given bcc material, that is a very different question. It can only be answered by looking for papers on the material in question.