Re: atomic position
dear everybody, I think there is some misconception here. It is true that the PDF-4+ is a commercial database: this means that sharing the information taken from the database (as recently done here) is an explicit violation of copyrights of ICDD and the violator might be legally prosecuted. When you use the database you save the time needed to look through the literature for the same information, paying the rights to the original owner and checking for errors in transcription and, sometimes, interpretation. That's why (at least) you pay for it. However, please update your records. The PDF-4+ is not the PDF-2. The PDF-4+ is a materials database that contains the "old" PDF-2 information, but it contains also structural data for a large quantity of compounds as well as properties and experimental patterns for several compounds. It contains also most of the ICSD coordinates that have been however thoroughly checked in house, quality marked, cross referenced and corrected when needed. So if the material is inorganic and the structure is not in PDF-4+, it usually means the structure is either very doubtful, too new or not yet extracted from the literature. In this particular case, as the chemistry and the space group are known, it is also likely that the compounds are isostructural to others already described in the literature Just notice that there are free alternatives to the commercial databases (e.g. COD, RRUFF) that in some cases can provide useful information, but in all cases the copyright warning exists Matteo (Chair of the BoD of the ICDD) - Prof. Matteo Leoni, PhD DICAM - University of Trento via Mesiano, 77 38123 Trento - Italy Tel +39 0461 282416 Fax +39 0461 282672 e-mail: matteo.le...@unitn.it 2016-06-14 9:04 GMT+02:00 Alan Hewat : > maybe they are in PDF4. >> > > PDF4 is a commercial database of d-spacings and unit cells derived from > them. It does not normally contain atomic positions, and not all structures > in PDF4 have been determined. ICSD does contain atomic positions if they > have been determined. > > Both PDF4 and ICSD are copyright and must not be posted to this mailing > list. As well, there is an explicit warning not to attach files to the > mailing list. Anyone who does so risks exclusion. > > There are other free databases that may contain the information you > require, but I doubt it. Try using Google Scholar to first determine if the > structures have been published. If not, you will have to determine them > yourself. It is not so difficult if you have good data and know the space > group and lattice dimensions. > > Alan > __ > * Dr Alan Hewat, NeutronOptics, Grenoble, FRANCE * > +33.476.98.41.68 > http://www.NeutronOptics.com/hewat > __ > > ++ > Please do NOT attach files to the whole list > > Send commands to eg: HELP as the subject with no body > text > The Rietveld_L list archive is on > http://www.mail-archive.com/rietveld_l@ill.fr/ > ++ > > > ++ Please do NOT attach files to the whole list Send commands to eg: HELP as the subject with no body text The Rietveld_L list archive is on http://www.mail-archive.com/rietveld_l@ill.fr/ ++
Re: PDF refinement pros and cons
several tricks to get it fast (histogram and distance binning/harvesting are among them, but there are several others). Again, focus on the problem: the narrower the peaks, the bigger the domains. Just guess the number of distances to be calculated for a domain of the order e.g. of 100nm. The number is large and you would need them all (together with their multiplicities) in order to calculate the pattern. You cannot use tricks: if you have defects you need to consider all distances (ok, well Monte Carlo can help there...) .. and we are simply skipping problems related to shape, size distribution, surfaces, texture, etc... cause the problem increases in computational demand. Application for now (in the powder diffraction world) is limited to the case of "well behaving" nano-sized powders The ideal solution: smart use of both worlds. If maths and physics agree on the data (reciprocal space) you can stop there. If there is still something missing, go for the real space methods. And check if the two are consistent cause in most cases it can be done with the right tools! M PS I am getting some of the other responses while I am writing.. so some info can be already outdated :oD In any case I see 3 players: pattern, PDF, structure The known choices are (if I'm not wrong): - get pattern from structure, compare pattern -> Rietveld - get pattern from random structure, compare pattern -> "Rietveld class" - get PDF from pattern, get PDF from structure, compare PDFs -> PDF refine - get PDF from pattern, guess structure from PDF -> PDF solve - get PDF from structure/microstructure, get pattern from PDF, compare pattern -> Debye with the condition that Rietveld is based on 3D periodicity, the PDF approach is not forced to! sounds reasonable? -- Matteo Leoni, PhD Department of Materials Engineering and Industrial Technologies University of Trento 38100 Mesiano (TN) Italy Tel +39 0461 882416e-mail:[EMAIL PROTECTED] Fax +39 0461 881977
Re: Line profile equation
Dear João I have read J. I. Langford's article (J. Appl. Cryst. 2000. 33, 964-974 ) and there I saw the following equation for the line profile (that was obtained from A.J.C. Wilson-1962): I(2theta) = [cos(theta)/lambda] I(s) I would like to know what the exact meaning of the cos(theta) factor in this equation is. What does it describe? I will be very grateful for some help. the term just account for the passage (done in the formula) from reciprocal space to 2theta space. Easy to demonstrate: s = 2 Sin[theta]/lambda so, taking the differential (for fixed lambda) ds = 2 Cos[theta]/lambda dtheta = Cos[theta]/lambda d(2theta) Enjoy M -- ---- Matteo Leoni, PhD Department of Materials Engineering and Industrial Technologies University of Trento 38050 Mesiano (TN) Italy Tel +39 0461 882416e-mail:[EMAIL PROTECTED] Fax +39 0461 881977Web:www.matteoleoni.ing.unitn.it
Re: Size Strain in GSAS
buna Nicolae, > Not only arithmetic, I think is clear that both and c were refined in a > whole pattern least square fitting. A private program, not a popular > Rietveld program because no one has inplemented the size profile caused by > the lognormal distribution. not sure no one did.. we're working with that kind of profiles at least since 2000 (published in 2001 Acta Cryst A57, 204), without the need for any approximation going through Voigts or Pseudo Voigts. Using FFT and some math tricks you can compute the "true" profile for a distribution of crystallites almost in the same time you calculate a Voigt curve, so why the need to use any approximate function? I think this agrees with what Alan just pointed out (well 5000 profiles per second if you do not include any hkl dependent broadening that has to be calculated for each of them (and perhaps for each subcomponent)... otherwise the speed reduces.. but yes few ms for each profile is the current speed for my WPPM code, implementing all this stuff within the WPPM frame). > > But the most important disadvantage is the necessity to choose the > > exact type of size distribution. For Sample 1 (which, obviously, have > > certain distribution with certain and c) you got quite different > > values of and c for lognorm and gamma models, but the values of Dv > > and Da were nearly the same. Don't you feel that Dv and Da values > > "contain" more reliable information about and c than those > > elaborate approximations described in chapter 6? > > Well, this is the general feature of the least square method. In the least > square you must firstly to choose a parametrised model for something that > you wish to fit. Do you know another posibility with the least square than > to priory choose the model? Without model is only the deconvolution, and > even there, if you wish a "stable" solution you must use a deconvolution > method that requres a "prior, starting model" (I presume you followed the > disertation of Nick Armstrong on this theme). also in this case it has ben shown possible to obtain a distribution without any prior information on its functional shape (J.Appl.Cryst (2004), 37, 629) and without taking the maxent treatment into account. I'm currently using without much problems for the analysis of nanostructured materials... advantages with respect to maxent are the speed and the fact that it can coexist with other broadening models (still not possible with maxent and still have to see a specimen where strain broadening is absent) and it's able to recover also a polydisperse distribution if it's present Just need to test it against maxent (if data would be kindly provided to do so). For the purists, just redo the calculation starting from different points and you can evaluate the error in the distribution using a MonteCarlo-like approach... As for the TCH-pV, well, it is no more than a pV with the Scherrer trend (1/cos) and the differential of Bragg's law (tan term) plugged in. This means it is ok as long as you consider a Williamson-Hall plot a good quantitative estimator for size and strain (IMHO). Mat PS I fully agree with Alan on the continuous request for Journals, but I bet the other Alan (the deus ex machina of the mailing list) should warn the members somehow... -- Matteo Leoni Department of Materials Engineering and Industrial Technologies University of Trento 38050 Mesiano (TN) ITALY
RE: Size Strain In GSAS
Leonid, > Could you, please, give a reference to a study where Dv and Da sizes > were derived from the parameters of pseudo-Voight or Voight fitted to > simulated profiles for various size distribution dispersions? I did something better (I hope).. at the end of the mesg you find xy data with a simulation: Ceria (Fm-3m), CuKa 0.15406 nm (delta function, i.e. no emission profile aberration)), size broadening due to a lognormal distribution of spheres only, no background, no noise, no Lorentz-Polarization, no aberrations of any kind. Peaks present in the pattern: 111 200 220 311 222 400 331 420 422 333/511 fitted but not used in the analysis as they have same d. Simulation done using WPPM, or, simply, taking the FT of the formulae proposed in Acta Cryst (2002) A58, 190-200 for the lognorm. The program used for the simulation allows fully recovery of the original parameters, starting from different initial values (self consistent check..). Analysis done using traditional "peak fitting" Williamson-Hall and Warren-Averbach methods, applying the formulae found e.g. on Phil Mag (1998) A77 [3], 621-640 to obtain the lognorm from the and values (should the ref be unavailable, they can be easily calculated, or I can also provide the formulae). I can send the simulation parameters and all plots/calculations I did to the interested members (just drop me a line). Results of WH and WA seems "good" but the distributions they provide are completely out with respect to the true one (unless I did some mistake, but that can be easily checked). I could have attached the results file here, but I'm sure Alan (as list moderator) would haven't been quite happy about attachments. Happy calculation, for those who wants to do the analysis by themselves without knowing the result in advance! Mat ----- Matteo Leoni, PhD Department of Materials Engineering and Industrial Technologies University of Trento 38050 Mesiano (TN) ITALY SIMULATED DATA in xy (2theta Intensity) format. 1.80e+001 4.470421e+000 1.806000e+001 4.513074e+000 1.812000e+001 4.556395e+000 1.818000e+001 4.600401e+000 1.824000e+001 4.645108e+000 1.83e+001 4.690529e+000 1.836000e+001 4.736682e+000 1.842000e+001 4.783659e+000 1.848000e+001 4.831332e+000 1.854000e+001 4.879787e+000 1.86e+001 4.929045e+000 1.866000e+001 4.979123e+000 1.872000e+001 5.030047e+000 1.878000e+001 5.081826e+000 1.884000e+001 5.134484e+000 1.89e+001 5.188043e+000 1.896000e+001 5.242526e+000 1.902000e+001 5.297955e+000 1.908000e+001 5.354352e+000 1.914000e+001 5.411743e+000 1.92e+001 5.470151e+000 1.926000e+001 5.529602e+000 1.932000e+001 5.590123e+000 1.938000e+001 5.651740e+000 1.944000e+001 5.714484e+000 1.95e+001 5.778379e+000 1.956000e+001 5.843577e+000 1.962000e+001 5.909879e+000 1.968000e+001 5.977428e+000 1.974000e+001 6.046257e+000 1.98e+001 6.116400e+000 1.986000e+001 6.187892e+000 1.992000e+001 6.260770e+000 1.998000e+001 6.335082e+000 2.004000e+001 6.410849e+000 2.01e+001 6.488118e+000 2.016000e+001 6.566933e+000 2.022000e+001 6.647338e+000 2.028000e+001 6.729377e+000 2.034000e+001 6.813098e+000 2.04e+001 6.898548e+000 2.046000e+001 6.985779e+000 2.052000e+001 7.074842e+000 2.058000e+001 7.165793e+000 2.064000e+001 7.258687e+000 2.07e+001 7.353585e+000 2.076000e+001 7.450544e+000 2.082000e+001 7.549629e+000 2.088000e+001 7.650906e+000 2.094000e+001 7.754684e+000 2.10e+001 7.860571e+000 2.106000e+001 7.968865e+000 2.112000e+001 8.079643e+000 2.118000e+001 8.192986e+000 2.124000e+001 8.308978e+000 2.13e+001 8.427707e+000 2.136000e+001 8.549279e+000 2.142000e+001 8.673760e+000 2.148000e+001 8.801265e+000 2.154000e+001 8.931897e+000 2.16e+001 9.065765e+000 2.166000e+001 9.202984e+000 2.172000e+001 9.343672e+000 2.178000e+001 9.487954e+000 2.184000e+001 9.635962e+000 2.19e+001 9.787833e+000 2.196000e+001 9.943709e+000 2.202000e+001 1.010374e+001 2.208000e+001 1.026809e+001 2.214000e+001 1.043692e+001 2.22e+001 1.061041e+001 2.226000e+001 1.078874e+001 2.232000e+001 1.097211e+001 2.238000e+001 1.116071e+001 2.244000e+001 1.135477e+001 2.25e+001 1.155511e+001 2.256000e+001 1.176082e+001 2.262000e+001 1.197270e+001 2.268000e+001 1.219102e+001 2.274000e+001 1.241607e+001 2.28e+001 1.264814e+001 2.286000e+001 1.288756e+001 2.292000e+001 1.313466e+001 2.298000e+001 1.338978e+001 2.304000e+001 1.365333e+001 2.31e+001 1.392571e+001 2.316000e+001 1.420732e+001 2.322000e+001 1.449861e+001 2.328000e+001 1.480009e+001 2.334000e+001 1.511227e+001 2.34e+001 1.543569e+001 2.346000e+001 1.577095e+001 2.352000e+001 1.611867e+001 2.358000e+001 1.647952e+001 2.364000e+001 1.685423e+001 2.37e+001 1.724355e+001 2.376000e+001 1.76
RE: Size Strain In GSAS
Leonid (and others) just my 2 cents to the whole story (as this is a long standing point of discussion: Davor correct me if I'm wrong, but this was also one of the key points in the latest size-strain meeting in Prague, right?) > Your recipe for estimating size distribution from the parameters of a > Voight-fitted profile is clear and straightforward, but I wonder have > you, or someone else, tested it on, say, simulated data for the model > of spherical crystallites having lognormal size distribution with > various dispersions? done several times... if you start from a pattern synthesised from a lognormal and you analyse it using a post-mortem LPA method (i.e. extract a width and a shape parameter and play with them to get some microstructural information), you obtain a result which (in most cases) does not allow you to reconstruct the original data (the Fourier transform of a Voigt and that of the function describing a lognormal distribution of spherical domains are different). I would invite all people using ANY "traditional" line profile analysis method to do always this check. Davor already pointed out cases where it works and cases where it does not: according to my experience those belonging to the first category are just a few. With a whole pattern approach and working directly with the profile arising from a distribution of domains, in most cases you're able to recostruct the original distribution without making any assumption on its functional shape (after all, most of the information to do so is contained in the whole pattern, even if it is well hidden). Concerning the Beyesian/maxent method, well, it is always a great idea, but unfortunately right now it is not mature enough to cope with a simple problem of combined instrumental, size AND strain broadening (unless something has been done in the last year). So ok it gives you the best result compatible with your hypotheses, but beware that "absence of any other source of broadening" should be listed among them.. and I'm not sure this is always the case! To put some water on the fire (otherwise it will burn all of us), I think the level of detail one needs on the microstructure, conditions the methods one's going to use to extract a result. No need to use highly sophisticated methods to roughly estimate a domain size (with an error up to +/- 50%) or to establish a trend within a homogeneous set of data, or also to obtain a better fit in the Rietveld method. Conversely, if a very high level of detail is sought, then I'd forget about a "traditional Rietveld refinement" and start approaching the problem from the microstructure point of view (after all, if one is interested in winning a F1 GP, he'd certainly not go for a Ferrari powered by a John Deere tractor engine!). cheers Mat - Matteo Leoni, PhD Department of Materials Engineering and Industrial Technologies University of Trento 38050 Mesiano (TN) ITALY
Re: rietveld refinement
just my 2 cents... > Could I be so stupid to say that such kind of works, including mine, are > nothing? following Nicolae, I should also add to the list myself as well as most people participating to the four editions of the size-strain conference/meeting/workshop and all participants to Davor's size-strain round-robin. I bet people should spend more time in the library... this is the point. This is also a self criticism as I'm not the best library addict (though, online resources has simplified life enormously)... We should not try to use line profile analysis methods as a black box: it is easy to obtain numbers from measured data (with a proper software a computer can do it automatically), but then it is in the ability of the scientist to attach them a proper physical meaning. What it is difficult (perhaps impossible?) is willing and pretending to do it in the general case as we're dealing with something that has no precise rules (domain size, shape and their distributions are not properties of the materials, nor they can be easily predicted in advance). Some simple cases have been studied and some references already posted by several people in here, and in most of them the agreement between diffraction and alternative techniques is quite good: just in few cases, though, enough information is available to interpret the strain broadening fully in terms of physical defects present in the material, or to model the size term using a more or less complex distribution of (iso-shape) domains. But also in those cases the result is the one compatible with the model assumptions and does not pretend to be "God's truth". So welcome the round robin on a more complex sample to test the maturity of the algorithms (they should be even tested on simpler examples, as concluded on the latest size-strain conference, but that's another story..), but beware that without any a priori info (or with a wrong one!), a vast set of odd results can be obtained. As a comparison, it would be like pretending to do a search match, a structure solution or, even worse, a Rietveld refinement on a material for which we don't know any chemical information... Going back to Leonid's question, well the answer is easy: check the premises... the assumptions behind the use of the TCH function are not compatible with he presence of a lognormal distribution of domains. It can be proven mathematically that the Fourier coefficients for a profile describing a lognormal distribution of domains have a hook at low Fourier number, hook that cannot be reproduced by any whatsoever voigtian or voigtian-like curve. This is a common problem in the use of Voigt and voigt-like curves in describing the peak profiles from nanocrystalline powders and is also the main source of the "superlorentzian" peak tails (they are a trick to get rid of the physical information contained in the profile ;) we are a bit masochist, aren't we?) Best regards Mat -- Matteo Leoni Department of Materials Engineering and Industrial Technologies University of Trento 38050 Mesiano (TN) ITALY Tel +39 0461 882416e-mail: [EMAIL PROTECTED] Fax +39 0461 881977Web: www.matteoleoni.ing.unitn.it
Re: GSAS informations
I just wanted to add my 2 cents to this argument... I think one big point in all the discussion on size and strain concerns the difference between what IS in the specimen, what we see with our probe (X-rays or neutrons, presumably) and what we reconstruct using A model. In most cases the model do not answer the easy question: "what is in the specimen?" I want to stress the "A" because in any case what we get is just a guess... subtle philosophers could speculate on this.. but I bet we're all scientists and not mere philosophers.. Anyway, people start their analysis with simple models and try to improve them.. easy but effective! As far as I know (well microelectronics is reality I think!) there is NO connection between grain shape (and therefore crystallite shape) and symmetry (or any descriptor for it). In this case the ellipsoid model could work, but is completely missing reality! The good scientist, however knows the limits of validity and the hypotheses on which the model is based (most modern scientists tend to forget this concept...) and knows that he obtained some "effective fit". Ok, specimens we analyse are simpler but... we are not dealing with specimens containing a set of perfect, equal, ordered, aligned crytallites.. More likely we have a distribution of shapes, sizes and orientation that can screw things up! In this case we can just hope that some simple model will accommodate all the mess! And using ANY anisotropic model is in most cases better than using none if you just care of "good fit". As for symmetry restrictions.. well.. they are welcome if they are consistent with the nature of diffraction (peak overlapping is always a painful problem in line profile analysis), but they are just related to our probe (X-rays, neutrones) and NOT with the original crystallites! Of course music changes if we talk about strain broadening. That's why Stephens models is good as an effective way of treating anisotropic broadening because "it fits better", but I'd not attach any physical meaning to the numbers you get out of it... I have written a bit too much I bet... so better if I go back to my size/strain modelling! OOps I was forgetting... Armel's replies to Nicolae emails are really great (you could be a good boxeur, Armel!), but I do not agree on one point: >Yes, anisotropic line broadening is rarely observed >with cubic compounds unless in very special cases >of faulting. this is true if you restrict the scope to size broadening only. Otherwise, line defects can be a source of anisotropic brodening in cubic materials (but you need a "good amount" to appreciate the effect)... Mat PS. There is no unique and simple solution to the problem... there are just scientists with their ability to attach the proper meaning to their results! -- w g( o 0 )g --oOO--(_)---OOo---oOO-w-OOo--- Department of Materials Engineering and Industrial Technologies University of Trento 38050 Mesiano (TN) Matteo Leoni, PhD ITALY .ooo0 0ooo. Tel +39 0461 882416e-mail: [EMAIL PROTECTED] ( ) ( ) Fax +39 0461 881977 \ (---) /-- \_) (_/ | | | | .ooo0 0ooo. ( ) ( ) \ ( ) / \_) (_/
Re: Rietveld method definition
Dear Armel (and all) > I just received the 48th Denver X-ray Conference > program, and was astonished by the title of the > D-092 conference : > "Ab initio structure solution as part of the Rietveld > refinement process" > > I believed till now that the Rietveld method was > the inevitable last part of an ab initio structure solution. > > Any hint ?;-). well.. this is already an old point of discussion with the guys in Bruker... the same "Ab-initio structure solution from powder data as part of the Rietveld refinement process" is also announced in their workshop in August in Glasgow. What is the boundary between "structure solution" and "Rietveld refinement"? It is true that there is a lot of confusion and the name "Rietveld" is used everywhere... but is it still a "Rietveld refinement" something that starts without actually knowing the structure? maybe it is an ab initio structure refinement... or... now we need more hints... > If the title of D-092 bears some truth, the SDPD mailing > list (http://www.cristal.org/sdpd/) should possibly > integrate the Rietveld mailing list... why not... but also the opposite would not be bad... ;o) Mat w g( o 0 )g --oOO--(_)---OOo---oOO-w-OOo--- MPI fuer Metallforschung Seestrasse, 92 74170 - Stuttgart (D) Matteo leoni, PhD Department of Materials Engineering University of Trento 38050 Mesiano (TN, Italy) .ooo0 0ooo.Tel +39 461 882417 ( ) ( )Fax +39 461 881977 E-mail: [EMAIL PROTECTED] \ (---) /-- \_) (_/ | | | | .ooo0 0ooo. ( ) ( ) \ ( ) / \_) (_/
Re: Stress
Hi all, ..following a bit the last posting on the topic... > << I start to be a little bit concerned about all those people claiming and > pointing out that diffraction don't measure a residual stress but a > residual strain. > .. > > So, my own idea is that people who don't know how to transform strain in > stress will measure only strain and the other are measuring stresses. > > Let me stress the audience about that; may be is that I am an engineer and > not a physicist or chemist.>> > > > Well, there is nothing wrong being an engineer but I cannot share your > concerns. Not to bore the audience, I just claim that you are just short of a > few quantities to measure stress directly from a diffraction experiment i.e. > of the elastic constants (which one hopes to find conveniently tabulated > somewhere and prays that they apply to the material under investigation). in any case it is in principle possible to measure some "elastic constants" using diffraction: it's just a matter of performing an in-situ 3 point bending or tensile or whatever kind of mechanical test on your sample and evaluate the response of the material itself. Since this is rarely possible, tabulated values as you say are the next choice... but they're useful only if you know what kind of grain interaction model is applicable to your particular case. I mean: you find tabulated the single-crystal elastic constants or the macroscopic mechanical elastic constants, but you need the X-ray elastic constants for the polycrystalline under study... you're not measuring all crystallites... you're just sampling some of them. Sometimes (eg the case of thin films, especially if they're textured), having the tabulated values for the single crystal elastic constants and getting a "physically reasonable" stress value are not synonyms! > For > measuring strain, on the other hand, the diffraction experiment provides all > the quantities necessary, you don't even have to know the material. are you really sure?? If you care about relative values I agree, otherwise you need the unstressed lattice spacing I agree on the fact that what you measure is something "easily" related to the interplanar spacing; from that you can "easily" get the integrated relative displacement in the laboratory system (with no "personal assumptions"). What follows is then just a matter of ASSUMING a behavior for the material and modelling. From my point of view it's a huge integral problem, so the solution is not unique... we just need better models As for the fact that someone is able to measure strain and some others are able to measure stresses, well, even if it is impossible to demonstrate that someone knows THE residual stress present in a sample, it's rather easy to prove that different values can be obtained from THE SAME set of measured points. Moreover, if variations of the sin2(psi) method are used, it can be shown even without recurring to Shannon's theorem, that in most cases the result depends also on the sampling in the (psi - 2theta/theta space) Ok ok... this is supposed to be the Rietveld mailing list ;o) oops.. Alan is gonna kill me... well.. but there are some analogies with the Rietveld method... and stress models have already implemented inside it... and well.. at least here neutron and X-ray fans cannot fight ;o) The two techniques are "more complementary each other" than for structure determination... to get more info (surface and bulk)... you need both! Now I'll get a lot of ... ehm on me.. but in this case really X-rays become a "surface" technique and neutrons a "bulk" one... otherwise I'd give anyone a thin film and a 30cm-thick "heavy" sample and ask to get reasonable strain values for both using only one technique! Have fun... Mat PS. In any case better to know the material.. and the instrument: residual macrostrain is not the only source for peak shift! w g( o 0 )g --oOO--(_)---OOo---oOO-w-OOo--- MPI fuer Metallforschung Seestrasse, 92 74170 - Stuttgart (D) Matteo leoni, PhD Department of Materials Engineering University of Trento 38050 Mesiano (TN, Italy) .ooo0 0ooo.Tel +39 461 882417 ( ) ( )Fax +39 461 881977E-mail: [EMAIL PROTECTED] \ (---) /-- \_) (_/ | | | | .ooo0 0ooo. ( ) ( ) \ ( ) / \_) (_/
