XRD Analysis - Science method

Hi Denis,

maybe look into FAULTS (part of the FullProf Suite). You can use it to simulate and refine layered structures with disordered layer sequence. You will need to define the different layers occurring in your material, the transition vectors from each layer to every other layer and the corresponding probabilities for each transition vector.

You will need to come up with a clever way to define the layers and transition vectors, but I think FAULTS should work here.

Cheers,

Jan

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Dear Achukwu,

I have mentioned all the ASTM standard about XRD analysis.

Hello,

From my experience, the XRD instrument usually comes with the software programs needed to analyze the data. For example, ours is a Bruker XRD that comes with “DIFFRAC” software to analyze the data. Our university has a license to this software which is how university users are able to access it. The software should allow you to find spectra library matches to identify your material and the degree to which your material matches the spectral fingerprint of another. After using the software tools, I typically verify the analysis by comparing across additional literature resources.

Complementing the Dalibor Matýsek answer, you can also try the ICSD database if you have access. The ICSD #36755 is very close to what you are looking for.

Obviously, Yes ! the EDM obtained by using FullProf suite software are acceptable for the publication. You can also follow my article.

Is there any one have icdd database pls?

Observation of Microstructure and Mechanical Properties in

Heat Affected Zone of As-Welded Carbon Steel by Using

Plasma MIG Welding Process

It was seen that the use of plasma MIG welding process has resulted in the refinement of the microstructure in the CGHAZ region, thus improving the mechanical properties of the as-welded SPCC steel. The highest microhardness values were obtained for conventional MIG in the CGHAZ regions. Incorporation of the plasma arc reduces the hardness, potentially increasing the ductility of the joining by plasma MIG welds. Further reduction in hardness can be obtained by decreasing the plasma current values.

Hi Jin Wook Yang,

I assume WQ and FC mean water quenched and furnace cooled, respectively. Is that correct?

My first idea was that the furnace cooled sample showed segregation of Ti and Zr, whereas the fast cooling of the WQ sample caused more uniform crystals.

It could also be that the transformation (bcc ht-Ti bcc -> hcp lt-Ti) in the WQ sample is (partially) suppressed due to the high cooling rate. This would cause that you observe both the reflections of the hcp and the bcc phase in you diffraction data. So the asymmetric reflections could actually be overlapping reflections.

Dear Usama Afzal . See the following useful link: https://web.iit.edu/sites/web/files/departments/academic-affairs/academic-resource-center/pdfs/Miller_Indices.pdf

First thing you definitely need is the best quality of xrd data. Both refinement tools you mentioned will always produce reliable outcomes, only if you thoroughly understand how refinement works.

Dear Derman Akşit . See the following useful links: https://www.rsc.org/cpd/teachers/content/filerepository/frg/pdf/Nanoparticles.pdf

As always, thank you Gustavo Henrique de Magalhães Gomes, for such a detailed information, and all other researchers, these answers are very insightful.

Single crystal diffraction is the gold standard for characterizing MOF(if you have single crystal)

Also check please the following useful link: https://tachantauclar.tistory.com/m/14

Shengquan Xuan you can obtain the structure factor from a CIF file using FullProf package software through Rietveld refinement analysis. You first need to perform a Rietveld refinement in FullProf and in the output configurations you can obtain the structure factors data from a GFourier software (included in FullProf). You can also perform an inverse Fourier transform in the Fourier density map obtained by GFourier to get accurate structure factors using mathematical software.

Yes, you can use the same sample for more techniques, but those techniques should not modify the sample, meaning that those techniques should not be invasive. In your case, you can use the same sample for UV-Vis, XRD, SEM, and FTIR.

Hi Abdul Moyez ,

XRD resolves structural information in 1D, a diffraction pattern on a detector resolves the structural information on a spherical/angular 2D slice. You would ideally need 3D scattering information (along the lines of CT) to be able to take an inverse Fourier transform of your data to exactly be able to identify your atomic positions. Once you regain your atomic position, you can carry out your desired MD simulation.

As a crude approximation, you can fill a 3D grid with the intensity obtained from your XRD pattern by assuming isotropy. That should allow you to at least generate a fictitious atomic arrangement that will exactly reproduce your XRD pattern. You can contact me for further help.

Dear Fabio, have you found out the reason and how to fix it. I am facing the exact same problem. Please help

Dear Unmesh Roy ,

please have a look at the attachment (Klug-Alexander tables 1-3)

or at the section 'Interplanar spacing' of:

Best regards

G.M.

Dear Sandip Hinge ,

it is the other way around;

compressive strain means decreasing atomic distances and lattice constants. Thus lattice plane distance(s) will also decrease a bit.

According to the Bragg-law^*) a decrease in d has to be compensated by a slight increase in theta; but not by lowering theta...

best regards

G.M.

*) lambda = 2* d*sin(theta)

Sir, how to find the peak values of XRD analysis, let me know

I can not give you an exact answer to what is happening. However, I try to answer what could be the reason for the difference between XRD and SAED patterns.

Considering a case of an ideal crystal, the XRD and SAED should give the same interplanar spacing. However, the difference in XRD and SAED patterns can be observed for non-ideal crystals (real samples).

The most apparent reason for the difference could be the fundamental difference in both techniques. SAED probes a very localized region of the sample ( what appears to be a couple of grain or polycrystalline agglomerates) under the microscope.

XRD on the powder or bulk samples, the information is from the coherent diffraction domains or grains, in large numbers (could be easily a billion grains), which includes grains on the surface area, illuminated by the x-ray beam, and some microns inside the surface ( penetration depth of x-ray, a few to 10s of microns, which depends on the material ). Thus, the XRD gives average information of large numbers of grains.

The local (SAED) and global (average information for the large sample volume; XRD) could differ in some cases. I could give you three main reasons and possible solutions to provide more structural information.

1-This could be the case where the sample has two phases (major- which appears in XRD; and minor appearing in SAED). The minor phase could be present in a small weight fraction. Therefore, the XRD signal from the minor phase will be weak and might not appear in XRD. However, if SAED is collected at the region where the minor phase is present (locally), peaks from the minor phase will appear in SAED.

To overcome this, You can collect high-quality XRD data (small steps, large collection time). The minor phase might appear. Alternatively, you can go the other way around and collect many SAED from different sample regions.

Please consider if the minor phase is in a very small weight fraction, below the detection limit (wight << 1%). You even might not see it in the high-quality XRD pattern.

2-It could also be the case that the minor phase is only present on the sample surface (obviously in small weight fraction); if the weight is small, the Bragg peaks for the minor phase will be very weak and might not appear in the XRD pattern, however, if you have performed the SAED only on the Surface region. You will see the minor phase.

To overcome this, you can perform Grazing incident XRD measurement, which is more sensitive to the sample surface. Alternatively, collect SAED from the surface and other regions.

3- It is also possible that some small weight fraction of the material might have preferred orientation or texture, which might not appear in XRD, but might appear in SAED. Especially if SAED is collected on the textured region of the sample.

Same as the last one, the texture will appear in GI-XRD.

If it is indeed (222) peak.

It could be a minor phase in small weight fraction, textured phase, or minor/textured phase on the surface region only. In these cases, it might not appear in XRD but appear in SAED.

Is there any free software for this?

XRD should show peaks for crystalline and amorphous phases as well. Loading may be low and not as reported (5%).

the cif is data from a structure already found, and you have to make the assumption; if yes then good

I think this article can help you.

New powder diffraction file (PDF-4) in relational database format: benefits and features of data mining

July 2002 Acta crystallographica. Section B, Structural Sciences 58 (Pt 3 Pt 1): 333-7

DOI: 10.1107/S0108768102002458

You can also find information on the web page

"Mineralogy Database"

Genelyn Rodriguez

Hello,

I see this as a classic problem, and that is that atoms in special positions cannot be refined. Which structural modification of BaTiO3 is to be refined here? For a cubic structure, it is not possible to modify any structural position.

Change the format of the CV.

and your CV is suitable for academics not for the Industrial Job.

Since most of your articles are under review you have to wait for a little bit.

Dear Preeti Sharma

you know, that the Scherrer*) equation deals with the crystallite size, but does not deal with with the particle size.

As far as you do not have got single crystalline particles, the shape of your particles (in your case 'rods') does not matter.

Just take the shape constant as about 0,9 and attach an error of about 20 % to 30% and that's it.

Please take in mind that the crystallite size value from the Scherr equation is a rough estimate of the mean of the crystallite sizes, but you sample exhibits a crystallite size distribution....

*) there is no 'Debye Scherrer equation'

Dear Gaurav Kumar ,

in addition to the answer already given above about ten years, pleas have a look at the answers of

https://www.researchgate.net/post/What-is-the-role-of-scan-speed-in-XRD-mineralogical-analysis-of-rock-samples.

With respect to your question there is nothing to add;

... but this thread here is very old one and only a very few of the RG colleagues will accidentally see your question here.

In general, one should file a new question here on RG. New questions will pop up very often and will stay some weeks in the 'Questions we think you can answer' section ...

Best regards

G.M.

Muralishwara Kakunje. Unfortunately, your diffraction patterns do not contain definable peaks, so you will not find d values for the nanoclay. There are two ways why this is so. Option 1 is that there was too little clay on the holder (and it was measured on the glass). The diffraction patterns then correspond to that glass. The second possibility is that the nanoclay is amorphous or nearly amorphous and does not provide measurable diffraction lines. I would see it in combination - little sample on the holder and it was measured on the glass.

The long distance disorder is essentially the same as the almost amorphous state, where order exists in the coordination sphere of the atoms themselves, but the resulting coordination polyhedra are stored chaotically.

Typically chi^2 is expected to decrease upon refining further parameters, if the refinement is well-behaved. As chi^2 increases, this suggests that the refinement is NOT well behaved at that point. This suggests high correlation of the refined paramaters push the refinement into a direction where everything is even poorer (you should see than in the visible quality of the fit, which should get poorer.). It is typically a sign for high correlation of the refined parameters, which can well be for uvw. This is in particular the case, if other parameters are being refined which do the same on the calculated profile as uvw. This depends, of course, on the details of you profile function and its refined parameters, and the history of your refinement. Based on this, and in view of the quality of your refinement do YOU have the feeling that you should touch u, v, w? I do not mean chi^2 (The way how people deal with their counting statistics frequently makes this number meaningless, unfortunately), but I refer to the agreement between observed and calculated profile!

Dear Meysam Shahrashoub

You did well using the modified hummers method to synthesize graphene oxide from graphite powder. In this respect, the graphite should used essentially of pure grade and not commercial grades. Now I recommend for you to review the procedure you followed with that in[1] under the sub item; 2.1 Preparation of graphene oxide by a modified hummer’s method ,pp471. Hope you find the missed step that preventing the graphite oxidation process progress.

Best regards

[1]https://pdf.sciencedirectassets.com/278653/1-s2.0-S1877705817X00167/1-s2.0-S1877705817316235/main.pdf

Dear Muhammad Faisal Iqbal

In the first procedure, check the presence of any other CuSe family member in your sample. It forms CuSe, Cu1.8Se, Cu3Se2, Cu2Se, in general, Cu1-xSe in a significant x range.

Due to the characteristics, the X-ray patterns for Cu1-xSe series with a slight difference in x will generate similar X-ray diffractions.

Another concern would be the preferred orientation, reducing it by pulverizing more than once the solid, which is very common in binary mixtures due to few stoichiometry variations.

If the synthetic route is ok, no preferred orientation is present, only one synthetic resulting material, and there is no reactant residue in the sample, you will be able to see only the indexed peaks of the desired phase.

Best regards,

WNM

Dear Khai Shenn Lau ,

in addition to the comments of Rajan Biswas and Dalibor Matýsek is not reasonable to exclude any peak or peaks without a proper identification.

Just as an example, gama-Fe2O3 if tetragonal, not cubic as found in specific synthetic route, may present some very low intensity peaks, easily identifiable as impurities or some other undesirable phase. However, those lines are very important for a proper Rietveld fitting, once they are superstructure lines, they just rise evidence of a stack problem on the tetragonal system. It happens in several other systems as well.

Before use the excluded region, you need to be sure those peaks don't belong to the present already identified phases.

Best regards

WNM

Rakibul Hassan

If it's the area under the peak that's of importance (because strain, size and instrumental broadening define the width) then it's clear that narrow peaks will have greater height (= more intensity) than wider peaks (= lower intensity). Narrow high intensity peaks are normally a feature of highly crystalline phases in appropriate alignment with the impinging x-ray beam.

Thank you Arabinda Nayak and Servando Chinchón Yepes for your answers. Those references were really helpful.

When calibrating I like to use the same material. Baring that, I want standards containing the same elements. I then optimize the instrument, including alignment. I then do 5 or 6 quant analyses of my standard, and average them together and use that final average quant as my standard.

For the general site you have only l = 2n as reflection conditions. The conditions you have listed are conditions for specfic, higher symmetry sites (if generated by spherically symmetric scatterers). However, your structure contains atoms on the general 12i sites. Hence the reflection condition is only l = 2n.

Thank you so much Dr. Gottschalk.

Origin pro 8.5.1

The thread that Sourav Roy recommended will get you the answer you need.

In summary of the thread: XRD measures crystallite size which is inherently different from particle size as measured by SEM. In general, particles are multi-crystalline (i.e. having multiple diffracting crystallites) which XRD inherently measures.

Good luck.

Hi Rakibul Hassan

Thank you for sharing the above equation.

I am not aware of this equation. To calculate the lattice parameter please use your XRD plot and compare it with the standard JCPDS data file to assign the (hkl) values. Once, you assigned the (hkl) values you can calculate the interplanar distance 'd' and then you the Bragg's equation to get the a, b, c values. You can get the crystallite size (D) by using Scherrer’s equation. For more information you may look at these: Current Catalysis 8 (1), 41-55, org/10.1016/j.jes.2019.09.004, 10.1021/acsomega.9b00252.

Regards

You need to generate a new set of equivalent xyz positions. This can be done by specifying the space group P4/nnc:2 instead of P4/nnc.

Hi Rakibul Hassan

The lattice constant, or lattice parameter, refers to the physical dimension of unit cells in a crystal lattice. Lattices in three dimensions generally have three lattice constants, referred to as a, b, and c. When a heteroatom with a smaller or larger radius was doped in a lattice structure e.g., C is getting doped in AB2O4 in place of B the lattice parameters of the doped material will change as compared to AB2O4. Thus, the crystallite size along with the unit cell volume will also change. So, by finding the correlation between the crystallite size and cell volume the reason can be explained from system to system. The above I can say from my side.

Regards

Area probed during XRR measurements depends on incident beam size. In horizontal direction, probed width is given by horizontal beam, whereas in vertical direction, beam footprint on the sample (=Vertical beam size/sin(theta)) defines the length probed during measurements. So, the length probed and hence, area probed during XRR vary with incident angle. For more information, you can check the article linked below:

It is preferable to use samples with size greater than beam footprint at critical angle of the material at incident energy. It can be done either by having bigger sample size or by decreasing vertical beam size using appropriate slits. In case both sample size and beam size can't be changed, you can go for measurements with knife edge so that only small region of the sample contributes to the reflected intensity. This is also given in the article linked above.

Regards,

How to calculate crystallite size from XRD by Scherrer equation in Origin #aminulcheminnovation

Dear Kamal Asadipakdel ,

I do not know the material of these domes.

It may be PMMA or polycarbonate or any other low atomic number material.

Low atomic number is essential for minimizing the x-ray attenuation.

As far as you do not use a spinning sample stage, there is no necessity to have a half sphere as 'dome' shape. In this case a half crylinder is sufficient (but with closed ends). In the cylindrical case you may use a thin kapton foil as wall of the half cylinder (tightened and glued to the circumference of the two half circle base plates). It is a little challenge of the design engineer(s) of your institution to set up such cap.

Good luck

It is not a good idea for one to introduce residual stress in the material. A residual stress is mainly related to when the processing of material is at a very high temperature. A residual stress can also be in the material when the material contained composition of atoms having different elements, i.e., atoms of more than one element. There are some other ways as well, which one can use to introduce a residual stress in a material. For that, you need to consult the preprint article given at the link. As suggested you in your earlier question, please refer the studies given at links http://dx.doi.org/10.13140/RG.2.2.27720.65287& https://doi.org/10.21203/rs.3.rs-88120/v2

The Best of Luck.

Rakibul Hassan,

Origin is always mentioned for this purpose. However, I find that the software, Fityk, is superior to Origin in the peak fitting department. I have provided the download link and an introduction video as well.

I find this software easier to use due to the flexibility, peak fit editing, and lite GUI as compared to the clunky peak fitting interface of Origin. The major plus is that you can place the peaks, run the fit, edit, and export all independently, where in Origin you don't get to do that as freely.

Good luck and cheers!

Intensity is an arbitrary unit, varies with instrument, source detector etc.and conditions during measurement as well. Please ensure you take readings from the same instrument if you want a comparative study. Or else go for relative intensity. Regarding higher intensity of a peak as compared to other, it defines the order of crystallinity in a certain plane of orientation.

Dear Hussein, many thanks for sharing this very interesting technical question with other RG members. Although we published at least 1,000+ crystals structures in the past ca. 40 years of chemical research, I'm absolutely not a proven expert in crystallography. However, to the best of my knowledge, there is no direct connection between hybridization and coordination number of a metal complex on the one hand and crystal system and space group on the other hand. Just take cobalt or nickel complexes like yours which are octahedrally coordinated with coordination number six. There are ten thousands or more complexes of this type in the chemical literature, and I'm sure that their structures encompass all crystal systems and space groups.

Also notable in this context is that completely analogous compounds can crystallize in different crystal systems. For a typical example (just one out of many) please have a look at the following interesting article:

(please see the attached pdf file)

I hope this helps answering your question. Good luck with your research and best wishes, Frank Edelmann

At 43 or 44 degrees of 2 theta LiMgBO3 and MgO appear

Did I understand it right as you say the Intensity of preferred orientation decreases at longer deposition? Does longer deposition really mean higher film thickness in your case? Maybe you are inducing strain during deposition? Are you sure that the observed reflex is the one from the preferred orientation? Maybe you can show some XRDs? I would assume the reflexes in preferred orientation to be broad at thin films and broaden going to thicker films.

Dear Heehyeon Sim ,

Humidity is responsible for broad PXRD as well as low crystalline phases or even amorphous.

Nevertheless, you report a compression step in the procedure. The compression contributes to the increase in the width of the peaks, as far as an alteration in the phase is observed.

Moving to higher two-theta is a sign of low lattice parameters as a result, that's for sure the phase is changing. The challenge is to identify the plane or planes where the phenomenon is affecting the material.

Indeed pair distribution function (PDF) is an experiment to set up under such conditions. A rocking curve experiment on the affected peaks also helps in understanding or defining the crystallographic directions of the effect.

Best regards,

WNM

Theoretically:

From XRD patterns as follows:

Variations in the height of the specimen in its holder, whether increasing or decreasing, cause XRD peaks to change from their original location.

Important and useful pathways are connected to the XRD Curve:

Scherrer’s equation:

Particle Size, D = (0.9 x λ)/ (d cosθ)

λ = 1.54060 Å (in the case of CuKa1) so, 0.9 x λ = 1.38654

Θ = 2θ/2 (in the example = 20/2)

d = the full width at half maximum intensity of the peak (in Rad) – you can calculate it using Origin software.

To convert from angle to rad

Rad = (22 x angle) / (7 x 180) = angle x 0.01746

Example: if d = 0.5 angle (θ)

= (22 x 0.5)/ (7x 180) = 0.00873 rad

Then

Delta=1/D^2

Dear Rajalakshmi Rajendran thank you for posting this interesting technical question. According to a common definition mentioned in the reference cited below "in zero-dimensional (0D) nanomaterials all dimensions are measured within the nanoscale (no dimensions are larger than 100 nm). Most commonly, 0D nanomaterials are nanoparticles." Thus it would be absolutely justified to call uniform nanoparticle of a size ranging around 40 nm in diameter zero dimensional (0D). For more information please have a look at this useful link:

For some more detailed information about synthesis and applications please also see the following very instructice review article:

(see attached pdf file)

I hope this helps answering your question. Good luck with your work and best wishes, Frank Edelmann

Nobody has an answer, so I'll write some observations. First of all, I would say that the phase analysis of clay minerals is one of the most complicated things that can be devised in mineralogy. Some clay minerals have virtually all the vices a mineral can have. Clay minerals show polytypism (all), often have a slightly or even strongly disordered or turbostratic structure, limited particle size and also display mixed layering too. All this affects not only (semi) quantitative but also qualitative analysis, and without a completely correct phase identification, quantitative analysis is meaningless.

But you can't say that it's not possible at all, it's just extremely complicated and you need to combine more techniques. The results of the Reynold Cup, which was organized by the Clay Minerals Society in the past, also prove that this is realistic.

Michael Harwarth

Thank you for your good explanation.

problem solved.

The reason for the data break was to open multiple data files at the same time. Thank you again.

You can easily get any XRD Reference or JCPDS file here,

Dear Shilpi,

I am from Xenocs, Xenocs is the leading market leader in SAXS/WAXS based in Grenoble -France and India.

In India, we have several installations few names already mentioned by respected persons. IIT- Bombay, RGIPT-Amethi, Lucknow etc.

If you need my help, I'll be very happy!

Suggested Readings

Good question.

The kernel average misorientation (KAM) is a measure for the local lattice curvature, applied to individual sets of orientation data points, which are typically obtained by diffraction experiments at microscopic resolution (such as in EBSD or TKD etc).

More specific the KAM quantifies the average misorientation around a measurement point with respect to a defined set of nearest OR nearest plus second-nearest neighbor points (u must double check also WHICH KAM value u have chosen). Values above a predefined threshold (often around 1 or 2 degrees) are usually excluded from the calculation, because these points are assumed to belong to adjacent grains or subgrains etc. Also check these threshold values before interpreting KAM values, maybe try out different threshold values.

This means that - strictly speaking - the KAM does not quantify the dislocation density at all but it serves as a local accumulative measure of averaged orientation changes in a small subset (of usually a larger orientation map) of orientation points relative to one centre orientation, valid for the specific lateral resolution of the underlying orientation point field, i.e. it is an orientation gradient measure.

It has been shown in the literature however that this gradient measure can indeed also serve as an approximation of a certain subset of the total dislocation content, namely as an approximation for those dislocations that act as the elementary carriers of the local lattice curvature (in other words local orientation change). These dislocations are called geometrically necessary dislocations following the original paper of Mike Ashby (M.F. Ashby, The deformation of plastically non-homogeneous materials, Philos. Mag. 21 (1970) 399–424. doi:10.1080/14786437008238426.).

This means that a high value of the KAM for instance derived for a given EBSD dataset is in the actual microstructure realised by a relatively high content of geometrically necessary dislocations (if not other misorientation carriers such as nano twins etc realise such an orientation change locally).

( One must however take care because sometimes high local KAM values can also come from poor surface quality of the sample / low-quality datasets. )

This also means - to the contrary - that very high statistically stored dislocation densities in microtexture arrays (which are devoid of mid-range orientation gradients) can in principle be present without producing a high KAM value signal.

Hence the answer to your question is that yes in principle you can have a high KAM value and hence a high density of geometrically necessary dislocations without necessarily having also a very high statistically stored density of dislocations and yes, the other way around is also possible, meaning that you can have a very high density of statistically stored dislocation densities at a low KAM value, namely when that array is curvature-free (no larger net misfit b-vector around that array).

Just remember that KAM represents curvature.

In principle for studying both quantities separately I would always recommend to combine KAM analysis with electron channeling contrast imaging, called ECCI, with which u can also probe the statistically stored dislocations in the scanning electron microscope.

Some further reading his here:

etc

enjoy

Dear Niranjan Haridas Menon!

In XRD, Peak intensity tells about the position of atoms within a lattice structure and peak width tells about crystallite size and lattice strain.

From your XRD plot, my suggestion is that whatever may be weight ratio w.r.t nanocomposite films with natural graphite filler, there will be slight variations in peak intensity value.

From the graph, the peak intensity at 3% loading is around 400. while considering peak intensity at 1% and 5% loading is around 800.

There is very little possibility of Peak intensity variations occurs while adding filler with a weight ratio of nanocomposite films.

I hope there may be a mistake/error in your reading at 3% loading. So plot the reading one more time and recheck.

I hope a Linear trend may be observed at 1%, 3%, 5%, and so on.

Best wishes.

The problem might be your baseline. The signal to noise ratio seems to be high hence the peaks might not be peaks but results of poor baseline. You might combine the data of multiple reading to improve this graph maybe.

Dear Kiran Dasari ,

there is equivalent CIF file for some COD database entry, which is free of charge, and the corresponding JCPDF proprietary database, now CCDC.

Please, check the Chrystallograohic Open Database (COD) website, download the free database and search by several ways, name, IUOAC, chemical formulae for the requested compound file.

Best regards

WNM

Variations in the height of the specimen in its holder, whether increasing or decreasing, cause XRD peaks to change from their original location.

Important and useful pathways are connected to the XRD Curve:

Scherrer’s equation:

Particle Size, D = (0.9 x λ)/ (d cosθ)

λ = 1.54060 Å (in the case of CuKa1) so, 0.9 x λ = 1.38654

Θ = 2θ/2 (in the example = 20/2)

d = the full width at half maximum intensity of the peak (in Rad) – you can calculate it using Origin software.

To convert from angle to rad

Rad = (22 x angle) / (7 x 180) = angle x 0.01746

Example: if d = 0.5 angle (θ)

= (22 x 0.5)/ (7x 180) = 0.00873 rad

Then

Delta=1/D^2

Hi dear Anas Rasras

MAUD software is using the Rietvel method to analysis and only you can do quantitative analysis with maud.

Williamson-Hall is another technique to analyze strain and crystallite size. You can do it easily with X'pert highscore plus software. Also, you can calculate easily FWHM with X'per highscore plus.

notice, These two methods are completely different from each other and you can't use the result of the Rietveld method to simulate the Williamson-hall plot.

By the way, maud is able to calculate phase fraction, crystallite size, strain, etc.

.

.

.If you have a problem drawing your Williamson-Hall plot, just send me your XRD data and tell me your material elements, phases, and fabrication process method.

Good luck

.

.

The shift in XRD peak of carbon materials is a very common phenomenon, which is largely observed in graphene, GO and rGO kinds of materials. a peak corresponds to 2θ≅12 °−17°, suggesting the inter-planar distance of GO get varied with the number of absorbed water molecules, with values from 0.61 nm for“dry” GO (2θ≅13 °−14°) to1.2 nm for“hydrated” GO (2θ≅8°−9°).

detailed explanation in the attached files.