The "state of the art" of diffraction analysis of crystallite size and lattice strain.
552 Z. Kristallogr. 223 (2008) 552–560 / DOI 10.1524/zkri.2008.1213 # by Oldenbourg Wissenschaftsverlag, Munchen ¨The “state of the art” of the diffraction analysis of crystallite size 1and lattice strainEric J. Mittemeijer* and Udo WelzelMax Planck Institute for Metals Research, Heisenbergstraße 3, 70569 Stuttgart, GermanyReceived June 1, 2008; accepted July 4, 2008X-ray powder diffraction / Line-profile analysis / many cases such information is not easily and statisticallyCrystallite size / Microstrain / Coherency of diffraction assured accessible by methods other than diffraction. The analysis of diffraction-line broadening, the topic ofAbstract. This paper addresses both old, but “renovated” this paper, evolved already shortly after the discovery ofmethods and new methods for diffraction line-profile diffraction of X-rays by crystals by Friedrich, Knippinganalysis. Classical and even extremely simple single-line and von Laue (1912): Scherrer (1918) found that themethods for separating “size” and “strain” broadening breadth of a diffraction line is related to the finite size ofeffects have merit for characterization of the material im- the diffracting crystals. Considering that, as follows fromperfectness, but it is generally very difficult to interpret differentiating Bragg’s law, lattice-parameter fluctuationsthe data obtained in terms of microstructure parameters as are also exhibited by diffraction-line broadening, Dehlin-used in materials science. Developments of recent years, ger and Kochendorfer, already as early as 1939, realized ¨focusing on distinct anisotropic line-broadening effects, as that a separation of the diffraction-line broadening in size-due to the type, orientation and distribution of dislocations and strain-related contributions can, in principle, beand minute compositional variation, will be touched upon. achieved provided that the diffraction angle dependence ofThe most promising development may be the synthesis of the line breadth is known .line profiles on the basis of a microstructure model and Unfortunately, straightforward extraction of quantitativeapplication of the (kinematical) diffraction theory without information on size and strain from the shape (“width”) dataany further assumption, which contrasts with the other is normally impossible. The least of the problems met ismethods. This approach can in principle be applied in sin- probably the elimination of instrumental broadening effects,gle-line and multiple-line variants and also in analyses of for which more or less reliable approaches exist whichthe whole diffraction pattern. The advantage is the direct depart either from recordings of broadening by standardevaluation of microstructure parameters as used in materi- specimens, or, more recently, from calculation of the instru-als science. The challenge is to develop microstructure mental broadening on the basis of the known instrumental/models which are flexible enough to be applicable in geometrical details of the diffraction experiment.more than one case . . . Fundamental problems are the unravelling of the var- ious contributions to the observed, broadened diffraction lines due to the various types of microstructural details, as crystallite size, lattice (micro)strain, planar faults (not con-1. Introduction sidered in this paper; cf. e.g. Refs. [3, 4]), etc., and their interpretation in terms of parameters that are used in mate-Diffraction lines of crystalline materials contain a wealth rials science, as dislocation densities, faulting probabilitiesof microstructural information: The amount and distribu- and crystallite size.tion of the phases in the material, compositional inhomo- On the one hand, more and more advanced methods togeneity, the crystallite size and shape distributions, the extract microstructural parameters from the profile para-crystallographic orientation distribution function, the con- meters of a single or multiple diffraction lines using morecentrations and distributions of crystal defects such as va- or less realistic, general assumptions on the material im-cancies, dislocations, stacking and twin faults, and, not perfection/line shape are developed: line-profile decompo-least, lattice distortions due to mechanical stresses, etc. sition. On the other hand, a recent, powerful, yet virginal(see, for example, Ref.  and references therein). In approach appears to be line-profile synthesis, where the microstructural parameters are determined by fitting line 1 Presented as keynote lecture at the International Conference on profiles, calculated on the basis of a model for the micro-the Diffraction Analysis of the Microstructure of Materials (“Size- structure specific for the material investigated, to measuredStrain V”; Garmisch-Partenkirchen, Germany, October 7–9, 2007).Full Proceedings available at www.zkristallogr.de (Z. Kristallogr. profiles (i.e. no line-shape assumptions are employed).Suppl. 27 (2008); open access). The length of the diffraction vector (and the correlation * Correspondence author (e-mail: firstname.lastname@example.org) of the positions of the scattering atoms) is crucial for the
The “state of the art” of the diffraction analysis of crystallite size and lattice strain 553occurrence of incoherency of diffraction, and as a conse- beam diffractometer is insensitive to defocusing errors.quence, apart from extreme cases, the so-called crystallite- This implies that the specimen can be tilted and rotated insize values depend on the reflection considered. Thereby, a parallel-beam diffractometer, as is required for stress andand this is not often realized, classical methods for line- texture measurements and investigations of the inhomo-profile analysis, but also recent developments, where all geneity and anisotropy of the microstructure , withoutreflections in the entire diffraction pattern are simulta- changing the (extent of) instrumental broadening. Thisneously analysed, under simple assumptions as a size (invariance of) instrumental broadening has been investi-broadening independent of the length of the diffraction gated both for diffractometers based on X-ray (polycapil-vector, can become invalidated. lary) lenses [12, 17] and X-ray mirrors [18, 19]. 2.2 Subtraction/incorporation2. Correction for instrumental broadening of the instrumental broadeningThe measurement apparatus, usually a diffractometer or Depending on the strategy of analysis of diffraction-linesome type of camera, generally brings about a significant broadening, line-profile decomposition versus line-profileintrinsic, instrumental broadening of the diffraction lines. synthesis (cf. Section 1), instrumental diffraction-lineTwo approaches can be considered to determine instru- broadening has to be taken into account differently.mental line broadening: On the one hand, a specimen with Whereas the former approach requires a subtraction of thenegligible structural line broadening can be investigated; instrumental from the measured broadening, the latter ap-the observed broadening is then (taken as) the instrumen- proach requires an ‘addition’ of the instrumental to thetal broadening. On the other hand, the broadening can be (calculated/modelled) specimen broadening. This is usual-calculated provided that sufficiently accurate models/meth- ly achieved by a deconvolution method and a convolutionods exist. The determination of the instrumental broaden- method, respectively.ing is dealt with in Section 2.1. Various approximate strategies for a correction of For analysing the line broadening measured for a parti- breadth parameters (as full widths at half maximum andcular specimen under investigation, a correction for instru- integral breadths) by “simple subtraction” of the corre-mental broadening has to be performed. This is dealt with sponding breadth parameters for the instrumental broaden-in Section 2.2. ing have been developed (cf. Refs. [20, 21] and references therein). Since the pioneering work of Stokes  invol-2.1 Determination of the instrumental profile ving a rigorous deconvolution of the measured broadeningThe selection of an appropriate reference specimen for with the instrumental broadening in Fourier space, no ma-characterising instrumental diffraction-line broadening re- jor progress in deconvolutive methods has been made untilquires careful consideration. In passing it is noted here recently: a novel method for deconvolution has been pro-that the standard reference material SRM 660 distributed posed both for laboratory  and synchrotron  pow-by the National Institute of Standard and Technology exhi- der diffractometers. This method combines scale transfor-bits a small, but in some cases, even for a laboratory pow- mation, data interpolation and fast Fourier transformationder diffractometer, non-negligible structural diffraction-line and permits a rigorous subtraction, in steps, of broadeningbroadening obscuring the genuine instrumental line broad- contributions due to axial divergence, flat specimen aberra-ening [5, 6]. A newer batch SRM 660a should lift this tion, specimen transparency and the wavelength distribu-problem (cf. also Ref. ). tion of the X-ray source, provided that the individual in- For a calculation of the instrumental diffraction-line strumental aberrations can be parameterized each by onlybroadening, the fundamental parameter [8, 9] and ray-tra- one diffraction-angle dependent width parameter.cing [10–12] approaches are usually adopted. Whereas theformer presupposes that the aberrations due to different in- 3. Size-strain broadeningstrumental aberrations can each be quantified by analyticalfunctions and can be treated independently, the latter is time 3.1 Simple approachesconsuming. An approach overcoming both drawbacks, byconsidering different aberrations simultaneously and provid- If data of high quality are unavailable (e.g. in the analysis ofing calculation efficiency, has been proposed recently for in-situ, non-ambient measurements) or an application doeslaboratory Bragg-Brentano powder diffractometers . A not merit the expenditure of time and effort required forcomparison of methods for modelling the effect of axial di- advanced line-profile analysis/synthesis methods (as wholevergence in laboratory powder diffraction arrived at the con- powder pattern modelling), a simple analysis of integralclusion that a computationally simplified approximation breadths may be appropriate for obtaining semi-quantitativebased on Edgeworth series can be employed . Special estimates of crystallite size and microstrain. Two basic ap-attention has also been paid to instrumental diffraction-line proaches for the separation of size and strain broadening onbroadening of synchtrotron-based diffractometers based on the basis of integral breadths can be adopted:collimating  and focusing  optics. (i) Single-line method . A diffraction line is con- In recent years diffractometers operating in parallel- ceived as a convolution of a Gaussian and a Lorentzianbeam geometry mode have become available also for la- (also called Cauchy) profile (i.e. as a Voigt function),boratory measurements [12, 16–19]. As the parallel beam where the Gaussian component is due to microstrain andgeometry does not rely on a focusing condition, a parallel- the Lorentzian component is due to finite crystallite size.
554 E. J. Mittemeijer and U. WelzelDetermination of the crystallite size (volume-weighed do- crystallite-size distribution and vice versa if a crystallitemain size in the direction parallel to the diffraction vector) shape is adopted; see e.g. Ref. ). In principle the col-D and of a measure for the width of the microstrain distri- umn-length distribution is given by the second derivativebution e is possible making use of the equations of the Fourier transform of the only size-broadened profile l [34, 35]. However, reliable determination of the column- bL ¼ ; ð1Þ length distribution on this basis suffers from problems due D cos q to background subtraction and truncation . In particu- bG ¼ 4e tan q ð2Þ lar, the obtained size distributions can be highly unreliablewhere l is the wavelength, 2q is the Bragg angle of reflec- in the presence of general strain broadening, which, in thetion, bL is the integral breadth of the Lorentzian compo- line-profile decomposition approach, has to be separatednent and bG is the integral breadth of the Gaussian com- from the size broadening on the basis of usually unvali-ponent (for details, see Ref.  and references therein). dated assumptions (corresponding results, for example ob-For the case of a Gaussian microstrain distribution it is tained in Ref. , should be mistrusted) .possible to calculate the local root-mean-square strain An alternative approach departs from the presupposi-h"2 i1=2 from e: h"2 i1=2 ¼ ð2=pÞ1=2 e [20, 26]. tion of a certain type of column-length or crystallite size 0 0 (ii) Williamson-Hall (WH) method . Assuming that distribution. For the description of monomodal distribu-the size and strain profile components are Lorentzian pro- tions, the Gamma- and lognormal distributions have beenfiles, the corresponding integral breadths are linearly addi- proposed:tive to obtain the total integral breadth in reciprocal space Gamma distribution :b* ¼ ðb cosqÞ=l (cf. Eqs. (1) and (2): 1 r pðnÞ ¼ n exp ðÀunt Þ ð4Þ 1 C b* ¼ þ 2ed * ; ð3Þ D where n denotes column length or crystallite size, C is awhere d * ¼ ð2 sin qÞ=l. A plot of b* versus d * should re- normalisation constant and r, u and t are adjustable para-sult in a straight line and the values for size and strain can meters (note that usually, t is (unnecessarily) taken asthen be obtained directly from the intercept and the slope one).of the straight line, respectively. Equation (3) presents one Log-normal distribution:specific expression for a ubiquitously adopted (but non-tri- pðnÞ ¼ ðð2pÞ1=2 sÞÀ1 exp ðÀðln n=no Þ2 =ð2s 2 ÞÞ=n ð5Þvial, see what follows in Section 3.5) assumption that sizebroadening does not depend on the length of the diffraction where no , the median, and s, the variance, are the adjusta-vector whereas strain broadening does. Other variants of ble parameters (cf., for example, Refs. [38, 39]). It hasthe WH method exist [e.g. adopting Gaussian shaped func- been found that in particular highly deformed metals oftentions, taking into account anisotropic line broadening (see exhibit log-normal column length/size distributions (e.g.Section 3.3.3), as due to dislocations etc.], but all are based Refs. [38–40]).on the assumption of specific profile shapes. Recently, the determination of column-length/crystal- Integral-breadth methods have been used in various also lite-size distributions without a prior assumption about therecent studies with the supposition that the results have a type of distribution has been attempted on the basis ofquantitative meaning (e.g. Refs. [5, 28–30]; for a critical whole powder pattern modelling. An approach involvingoverview of such methods, see also Ref. ). Results quan- histograms with “tuned” bin width and adjustable bintitatively consistent with results obtained by more advanced height, but assuming a spherical crystallite shape, has beenmethods can be obtained, in particular and obviously for proposed in Ref.  (see Fig. 1).cases where one source of line broadening prevails [31, 32]. An approximate solution for obtaining the crystallite- size distribution together with information on crystallite3.2 Column length/crystallite size distribution shape has been proposed in Refs. [42, 43]. In the latter ap- proach a microstrain distribution with a homogeneous strainThe column length will generally exhibit a distribution in each crystallite has been presupposed (this can be a se-(the column-length distribution can be calculated from the vere limitation; cf. Section 3.5 and see next paragraph).Fig. 1. Whole powder pattern modelling: Crystallite size, D (diameter of the sphere (¼ crystallite)), distributions, pðDÞ, of nanocrystalline ceriapowders calcinated for 1 hour at different temperatures (increasing from the left to the right). The full histogram is the result of the analysiswithout prior assumption on the crystallite size distributions, whereas the line is the result of the analysis restricted to a log-normal size distribu-tion. Taken from Ref. .
The “state of the art” of the diffraction analysis of crystallite size and lattice strain 555 The two categories of approaches (i.e. whether or not the distance between the defects (projected onto the dif-assuming a distribution function) for determining the col- fraction vector), the probability function for the amplitudeumn-length distribution (and, possibly, the crystallite-size of the component strain fields and a function describingdistribution on the basis of an additional assumption on the average shape (width) of the component strain fields.crystallite shape) both require that broadening from In the simplest case for application of the strain-field mod-sources other than finite size (as microstrains) is marginal el, the Fourier coefficient for the only strain broadenedor absent: the required assumptions to separate the size profile Ad ðLÞ is described by only three parameters (cf.broadening from the other broadening components renders Eq. (7) of van Berkum et al. ): (i) the mean projecteda subsequent determination of column-length or crystallite- (onto the diffraction vector) defect distance hsi, (ii) thesize distributions unreliable (corresponding results, e.g. as root-mean-square strain he2 i, and (iii) the width of the opublished in Ref.  (see first paragraph of this section) (Lorentz shaped) component strain fields, w. A componentor Ref.  (see above paragraph) should be mistrusted). representing a possible size broadening can simply be in- cluded . For applications of this strain-field model to3.3 Microstrain broadening ball-milled metal powders, see Refs. [39, 47, 48]. Methods departing from specific microstructural modelsWhereas the fundamentals of size broadening are well es- have been developed for analysing line broadening due totablished and in a mature state already since the 1950s, inclusions in a crystalline matrix  and due to disloca-[34, 35] analysis of strain broadening is a field of currently tions. In the following the focus is on dislocation linestrong activity, where both methods imposing assumptions broadening.on the kinematical diffraction theory and methods depart- The pioneering work in this field is due to Krivoglazing from a microstructural model are developed. and Ryaboshapka  and Wilkens . Krivoglaz and Ryaboshapka considered sets of statistically random distri-3.3.1 Methods imposing assumptions butions of non-interacting (edge or) screw dislocations. on the kinematical diffraction theory Wilkens demonstrated that a random distribution of dislo- cations (in a set) is unrealistic and introduced the conceptAn overview of the methods based on specific assump- of the restrictedly random dislocation arrangement. To thistions about the strain distributions in materials without re- end, the degree of correlation in the dislocation distribu-ferring to a specific microstructural model is provided by tion of a set was described by the so-called cut-off radiusTable 1 [27, 35, 44–46]. Re , which can be considered as the radius of a cylinder A quantitative evaluation of size and strain parameters within which the dislocation arrangement is random: Noderived from broadened line profiles requires thorough con- elastic interaction of the various dislocations sets in thesideration of the underlying assumptions in the methods crystal is considered to occur.used. A comparative application of the different methods to The strain Fourier coefficients Ad ðLÞ can be approxi-an imperfect material is not straightforward because the as- mated by :sumptions are incompatible (e.g. Gaussian strain distribu-tions in the Warren-Averbach method versus small strain gra- Ad ðLÞ ¼ exp ½ÀðcLÞp ; ð6Þdients in the alternative method) and the resulting parameters d where c characterizes the width of A ðLÞ and the exponentare not defined in the same way (e.g. volume-versus area- p takes values between 1 (Lorentzian line profile) and 2weighted crystallite size) ; see also Fig. 1 in Ref. . (Gaussian line profile). The shape parameter M,3.3.2 Methods departing from a microstructural model M ¼ Re ðrÞ1=2 ; ð7ÞA flexible general method based on a microstructural can be calculated from p (cf. Eq. (2.19) of Vermeulen et al.model without referring to a particular type of defect is ). c is related to the square-root of the dislocation densitythe strain-field model proposed by Van Berkum et al. . r. Note that dislocation line broadening is usually anisotro-In this approach, the strain field is composed of a super- pic, i.e. it depends on the hkl reflection (i.e. it depends on theposition of the (component) strain fields of individual de- orientation and length of the diffraction vector; cf. Sectionfects. The strain fields of the lattice defects are described 3.3.3). This can be rationalized by the so-called dislocation-statistically by three functions: the probability function for contrast factor, which is contained in c in Eq. (6).Table 1. Summary of basic assumptions made in line profile decomposition methods and the type of size and strain data obtained . Ad ðLÞ isthe strain (‘distortion’) Fourier coefficient of a line profile, L is the correlation distance perpendicular to the diffracting planes.Method Assumptions Size StrainWilliamson-Hall conventional Lorentz shaped peak Volume-weighted Maximum strain, e related to localplot [27, 44] 1949, 1953 profiles for size- and strain- column length mean squared strain he02i broadened profile for Gaussian strain distributionsWarren-Averbach [35, 45] Gaussian strain distribution Area-weighted column length Mean squared strain, he2(L)i,1950, 1952 or small strains related to Ad(L)Alternative method  1994 Small strain gradients Area-weighted column length No analytical relation between and broad size distribution he2(L)i and the strain Fourier coefficients
556 E. J. Mittemeijer and U. Welzel Dislocation densities and configurations have been in- tropic diffraction-line broadening may be categorized asvestigated in thin films and plastically deformed materials follows:(see, for example, Refs. [39, 54, 55]). In addition to the (i) Only small (negligible) microstrain gradients withindislocation density and the cut-off radius, the fractions of crystallites. In this case, the increase of line broadeningscrew and edge dislocations can be determined. For a re- with increasing length of the diffraction vector, for a givencent review on dislocation line broadening, see Ref. . set of diffracting lattice planes, is proportional to tan q (cf. also Section 3.5) [47, 57]. Phenomenological models for3.3.3 Anisotropic microstrain-like diffraction-line this type of anisotropic microstrain diffraction-line broaden- broadening ing have been developed and implemented in Rietveld-re- finement programs (see, for example, Ref. ). A modelThe occurrence of anisotropic diffraction-line broadening case for this type of line broadening is a (hypothetical) iso-(i.e. the diffraction-line broadening depends non-monoto- tropic microstress distribution which, in combination withnously on the hkl reflection when plotted versus 2q) is a single-crystal elastic anisotropy, results in an anisotropicquite general phenomenon which has attracted consider- microstrain distribution .able attention both in phenomenological and microstruc- This approach is likely to overestimate the anisotropy ofture-based modelling of diffraction-line broadening. Aniso- diffraction-line broadening, as an isotropic microstress distri- bution gives rise to geometrically incompatible strains in dif- ferently oriented crystallites. The real grain interaction in a polycrystalline material is more likely to be between isotro- pic stress and isotropic strain distributions. Another recently considered source of anisotropic (microstrain-like) line broadening are composition fluctuations in a non-cubic ma- a a bFig. 2. (a) Full width at half-maximum (FWHM) of the reflections ofa e-FeN0.433 powder and LaB6 (used for the determination of the in-strumental line broadening) measured using a Bragg-Brentano dif- bfractometer with Co Ka radiation. The apparent ‘scatter’ of the linewidths of the powder is due to compositional inhomogeneities. Fig. Fig. 3. (a) The FWHM (w*) and the integral breaths (B*) as a func- f2a and b have been taken from Ref. . (b). The anisotropy of the tion of the reciprocal space coordinate d* in the classical Williamson-microstrain-like broadening observed from a e-FeN0.433 powder. The Hall plot in the case of Nb ball milled for one day. The indices of thedirection dependence of the FWHM, BfD2 q;hkl , as a function of the reflections have also been indicated in the figure. Note the pronounc-angle of the diffraction vector relative to the c axis for the hexagonal edly anisotropic nature of line broadening. (b) The modified William-crystals system. The separate points indicate the experimental data; son-Hall plot of the same data as in Fig. 3a. As a function of d *C 1/2the solid line represents the curve obtained by fitting a model for line (where C is the dislocation contrast factor). The indices of the reflec-broadening due to compositional fluctuations to the experimental tions have also been indicated in the figure. Fig. 3a and b have beendata. A compositional fluctuation of e-FeN0.433Æ0.008 is obtained. taken from Ref. .
The “state of the art” of the diffraction analysis of crystallite size and lattice strain 557terial (see Fig. 2 for an example) . For a recent general a so-called columnar microstructure occurs, where the filmtreatment on anisotropic microstrain broadening due to a consists of e.g. columnar-shaped grains separated by grainfield-tensor (rank 0, pertaining to composition variation; rank boundaries oriented more or less perpendicularly to the2, pertaining to stress/strain distributions), cf. Ref. . layer surface. For such a thin film, the crystallite size is an (ii) No assumption about microstrain gradients: Adop- anisotropic quantity: the crystallite size along the surfacetion of a microstructural model. In this case, the depen- normal is much larger than the crystallite size in the planedence of strain broadening on the length of the diffraction of the film. Thus, macroscopically anisotropic size broad-vector follows from the microstructural model. The re- ening occurs (see Fig. 4a for an example) . Anisotro-cently most frequently studied case is dislocation line pic size broadening can be accompanied by anisotropicbroadening, for which anisotropic line broadening is due strain broadening (see Fig. 4b) , which can also occurto the orientation of the diffraction vector with respect to due to unequal densities of defects (as dislocations on dis-slip systems and the anisotropy of elastic constants (cf. tinct slip systems) along different directions in the speci-Section 3.3.2 and Fig. 3; e.g. Refs. [56, 61–63]). men . The analysis of macroscopically anisotropic diffraction-3.4 Macroscopic anisotropy line broadening is considerably simplified experimentally by the use of parallel-beam diffractometers, because instru-Massive and polycrystalline specimens generally exhibit mental aberrations occurring for focusing diffractometersan anisotropic microstructure. Consider, as an example, (i.e. ‘defocusing’) upon changing the orientation of thethin films deposited by physical vapour deposition: Often diffraction vector (from e.g. the specimen surface normal direction, for the case of Bragg-Brentano diffractometers) can be avoided . 3.5 Crystallite size and coherency of diffraction For most polycrystalline specimens, the phase difference (reduced modulo 2p) of a wave scattered by one crystal- lite and the wave scattered by a second crystallite takes values between 0 and 2p with equal probability. In this case, the total diffracted intensity can be taken as the sum of intensities scattered by the individual crystallites sepa- rately. This naturally leads to the usually adopted concept of size broadening due to the finite size of individual crys- tallites and strain broadening related to the relative displa- a cement of atoms within one grain. A more general ap- proach is to consider the whole irradiated volume of a polycrystal as a coherently scattering domain. Such an ap- proach has been followed by van Berkum et al. for analys- ing strain broadening on the basis of a flexible model for strain fields associated with lattice defects (cf. also Section 3.3.2) . As the phase difference of scattered waves originating from different scatterers (atoms) is the scalar product of the diffraction vector and the position (differ- ence) vector of the scatterers, both the character of the strain fields in a specimen and the length of the diffraction vector are decisive for diffraction-line broadening. It has been demonstrated that for general strain broadening, the order-dependence of the diffraction-line width is complex, b i.e. neither order-independent broadening (traditionally termed ‘size broadening’) nor broadening proportional toFig. 4. (a) The crystallite, grain sizes of a 250 nm thick Ti3Al layer the length of the diffraction vector (traditionally termedas viewed along different hhkli* directions, i.e. as function of the ‘strain broadening’) occurs (see Fig. 5a). Two limitingtilting angle w. The schematic figure represents the rectangular Ti3Algrains (with a height of 50 nm and a width of 6 nm) in the Ti3Al cases have been identified:layer, the crystallite size of the rectangular grain DhhkliÃ measured (i) For infinitely broad component strain fields of lat-along the hhkli* direction is 50 nm/cos w for 0 w 6.89 ; 6 nm/ tice defects (w=hsi ! 1; cf. Section 3.3.2), the broaden-sin w for 6.89 w 90 , as shown by the solid and dashed lines. ing is proportional to the length of the diffraction vector(b) The microstrains and the grain boundary fraction FGB of the d * (cf. Eq. (3)):Ti3Al layer as viewed along different hhkli* directions, i.e. as func-tion of the tilting angle w. FGB ¼ gD=D, where D is a constant rela-tive to the grain boundary thickness, and equals 1.0 nm here; g is a b ¼ ð2pÞ1=2 d *he2 i1=2 ; ð8Þgeometrical constant and equals 1. The results suggest that FGB andthe microstrain behave similarly as function of w. Fig. 4a and b have where he2 i is the mean squared strain. This is the well-been taken from Ref. . known strain broadening for a specimen with a constant
558 E. J. Mittemeijer and U. Welzel and constant at large lengths; see Fig. 5b) has recently been experimentally confirmed on the basis of measured line broadening of nanocrystalline thin films by Rafaja et al. (see Fig. 5b) : i.e. a loss of coherency with in- creasing length of the diffraction vector at small lengths of d * leading to incoherent diffraction of the individual crys- tallites at larger length d*. For another study of the effect ´ of partial coherence on size broadening, see Ribarik et al.  For an application of the strain-field model incorpor- ating coherency effects, see Lucks et al. . a 4. Concluding remarks and perspectives (1) Whereas unprejudiced individual peak-profile analysis (still) allows the most severe microstructural model testing, a simultaneous analysis of all reflections, sub- ject to more severe (e.g. profile-shape) constraints, is required if overlap of diffraction lines occurs. Distinc- tion of both evaluation approaches can be expected to diminish as more sound microstructure-based diffrac- tion models are developed. (2) For sensitive crystal-imperfection analysis, an evalua- tion of anisotropic diffraction-line broadening, with re- spect to the (orientation and) length of the diffraction b vector in either the reciprocal lattice or the specimen frame of reference, is a prerequisite.Fig. 5. (a) Application of the strain-field model: Relative integral (3) Preference for microstructure model-based line-profilebreadths of line profiles in reciprocal space for different relative synthesis methods over line-profile decompositionwidths w, of the component strain fields. The subscript ‘r’ indicates methods is due to the application of flawed micro-that the respective quantity has been normalized by the mean pro-jected (on the diffraction vector) defect distance. Taken from Ref. structure/diffraction models in line-profile decomposi-. (b) Diffraction line broadening measured for a nanocrystalline tion methods.Ti0.38Al0.62N thin film (points) and the numerical simulation (solid (4) Is a general, practically applicable microstructure/dif-line) taking into account the partial coherence of neighbouring nano- fraction model possible?sized crystallites. Taken from Ref. . Note the similarity with re-sults shown in Fig. 5a for, e.g., wr ¼ 0:03. Referenceslattice spacing d within each crystallite and a Gaussian  Mittemeijer, E. J.; Scardi, P. (Eds.): Diffraction Analysis of thespacing distribution over the crystallites, where: Microstructure of Materials. Springer, Germany 2004.  Dehlinger, U.; Kochendorfer, A.: Linienverbreiterung von ver- ¨ he2 i ¼ ðhd2 i À hdi2 Þ=hdi2 : ð9Þ formten Metallen. Z. Kristallogr. 101 (1939) 134–148.  Estevez-Rams, E.; Penton Madrigal, A.; Scardi, P.; Leoni, M.:This proportionality of line broadening (e.g. integral Powder diffraction characterization of stacking disorder. Z. Kris- tallogr. Suppl. 26 (2007) 99–104.breadth) and length of the diffraction vector (see Fig. 5a)  Leoni, M.: Diffraction analysis of layer disorder. Z. Kristallogr.has often been used for strain broadening in general (cf. 223 (2008) 561–568.Eqs. (2) and (3); Section 3.3.1).  Pratapa, S.; O’Connor, B.; Hunter, B.: A comparative study of (ii) For infinitely narrow component strain fields single-line and Rietveld strain–size evaluation procedures using MgO ceramics. J. Appl. Cryst. 35 (2002) 155–162.(w=hsi ! 0, as corresponds to small-angle grain bound-  Masson, O.; Dooryhee, E.; Fitch, A. N.: Instrument line-profile ´aries in polycrystalline material; cf. Section 3.3.2), it fol- synthesis in high-resolution synchrotron powder diffraction. J.lows that Appl. Cryst. 36 (2003) 286–294.  Balzar, D.; Audebrand, N.; Daymond, M. R.; Fitch, A; Hewat, b / ðd *Þ2 ð10Þ A.; Langford, J. I.; Le Bail, A.; Louer, D.; Masson, O.; McCo- ¨ wan, C. N.; Popa, N. C.; Stephens, P. W.; Toby, B. H.: Size-for small lengths of the diffraction vector. For larger lengths, strain line-broadening analysis of the ceria Round-Robin sam-a constant line width b is obtained (‘classical’ size broaden- ple. J. Appl. Cryst. 37 (2004) 911–924.  Cheary, R. W.; Coelho, A. A.: A fundamental parameters approaching). A model microstructure producing this type of broaden- to X-ray line profile fitting. J. Appl. Cryst. 25 (1992) 109–121.ing consists of domains of undistorted material shifted with  Kern, A.; Coelho, A. A.; Cheary, R. W.: Convolution based pro-respect to each other. For increasing length of the diffraction file fitting. In: Diffraction Analysis of the Microstructure of Ma-vector, the phase differences (modulo 2p) are almost uni- terials (Eds. E. J. Mittemeijer, P. Scardi) pp. 17–50. Springer, Germany 2004.formly distributed, thus incoherency of diffraction occurs.  www.bgmn.de This prediction of this type of diffraction-line broaden-  Bergmann, J.; Friedel, P.; Kleeberg, R.: BGMN – A new funda-ing (increasing at small lengths of the diffraction vector mental parameters based Rietfeld program for laboratory X-ray
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560 E. J. Mittemeijer and U. Welzel Leineweber, A.; Mittemeijer, E. J.: Diffraction line broadening  Dragomir, I. C.; Ungar, T.: The dislocations contrast factors of ´ due to lattice-parameter variations caused by a spatially varying cubic crystals in the Zener constant range between zero and scalar variable: Its orientation dependence caused by locally unity. Powder Diffraction 17 (2002) 104–111. varying nitrogen content in epsilon-FeN0.433. J. Appl. Cryst. 37  Leoni, M.; Martinez-Garcia, J.; Scardi, P.: Dislocation effects in (2004) 123–135. powder diffraction. J. Appl. Cryst. 40 (2007) 719–724. Leineweber, A.: Anisotropic microstrain broadening due to field-  Rafaja, D.; Klemm, V.; Schreiber, G.; Knapp, M.; Kuzel, R.: tensor distributions. J. Appl. Cryst. 40 (2007) 362–370. Interference phenomena observed by X-ray diffraction in nano- Ungar, T.; Gubicza, J.; Ribarik, G.; Borbely, A.: Crystallite size ´ ´ ´ crystalline thin films. J. Appl. Cryst. 37 (2004) 613–620. distribution and dislocation structure determined by diffraction  Ribarik, G.; Audebrand, N.; Palancher, H.; Ungar, T.; Louer, D.: ´ ´ ¨ profile analysis: Principles and practical application to cubic and Dislocation densities and crystallite size distribution in nanocrys- hexagonal crystals. J. Appl. Cryst. 34 (2001) 298–310. talline ball-milled fluorides MF2 (M ¼ Ca, Sr, Ba and Cd), deter- mined by X-ray diffraction line-profile analysis. J. Appl. Cryst. 38 (2005) 912–926. Zeitschrift für Kristallographie New Supplements to be Published International journal for structural, physical, Z. Kristallogr. Suppl. 27 (2008) and chemical aspects of crystalline materials Proceedings of the ”5th Size Strain“ Conference ZEITSCHRIFT FÜR KRISTALLOGRAPHIE (Diffraction Analysis of the Microstructure of Materials) held in October 2007 in Garmisch-Partenkirchen, Germany 16. Jahrestagung Editors: Eric J. Mittemeijer, Paolo Scardi, Andreas der Deutschen Gesellschaft für Kristallographie f Erlangen, 3.– 6. März 2008 Leineweber and Udo Welzel Referate Oldenbourg Z. Kristallogr. Suppl. 28 (2008) Supplement Issue No. 28 2008 Abstract of the 16th Annual Meeting of the German Crystallographic Society Oldenbourg held in March 2008 in Erlangen, Germany www.zkristallogr.de