JP6645809B2 - Microscopic analysis method of fine particles and analyzer used for microscopic analysis method of fine particles - Google Patents

Description

  The present invention relates to a microscopic analysis method for fine particles and an analyzer used for the microscopic analysis method for fine particles.

  Since the industrial revolution, for example, with respect to materials such as metals, research and development utilizing characteristics of particles having a size on the order of micrometers or more have been carried out, and have contributed to the development of society.

In addition, semiconductor materials that have rapidly developed after the war have constituted circuits having more than 100 million transistors, and have produced products such as supercomputers, liquid crystal televisions, and smartphones.
The size of the semiconductor material is on the order of nanometers in the thickness direction, that is, in the direction perpendicular to the substrate, but is on the order of centimeters in the in-plane direction of the substrate.

  When measuring the atomic position of a material such as the above-mentioned metal or the like, or measuring the atomic position of a semiconductor material in the in-plane direction, the size of the above-mentioned material is smaller than the interatomic distance to be measured, that is, in the order of angstrom. Therefore, the material can be measured assuming that the material has an infinite periodicity in which substantially infinite atoms are arranged.

  In recent years, research and development have been actively conducted on materials having a size that cannot be regarded as having an infinite number of atoms arranged in a row and inevitably have a finite periodicity.

The material having the finite periodicity, a material having a finite periodicity, such as nanoparticles, because of its size, characteristics that are incomparable with the material that can be considered to have the infinite periodicity, They may exhibit completely different properties.

  As such nanoparticles, for example, a surface reaction peculiar to nanoparticles occurs by irradiation with light, a photocatalyst capable of purifying harmful organic compounds and fine substances, a fuel cell catalyst having a low environmental load, and the like. Materials that are indispensable for social development in the 21st century can be mentioned.

  Such a nanoparticle cannot be regarded as having an infinite periodicity, but can be regarded as having a periodicity within the particle, that is, within a region on the order of nanometers.

  As a method of performing observation and analysis in a region on the order of nanometers, there are (1) an observation and analysis method that can be performed at an atomic level, and (2) an observation and analysis method that can capture finite periodicity.

  Non-Patent Documents 1 and 2 disclose techniques corresponding to the observation and analysis methods that can be performed at the atomic level in the above (1).

In Non-Patent Document 1, an aberration-corrected High-Angle Annular Dark Field Scanning Transmission Electron Microscopy (HAADF STEM) is used as an observation unit.
The HAADF STEM observation is performed under a low accelerating voltage, and the obtained image is calibrated by simulation, thereby achieving a three-dimensional image of the nanoparticles at a level at which atoms can be clearly seen.

In Non-Patent Document 2, Scanning Transmission Electron Microscopy (STEM) is used.
Then, the rotation axis of the STEM is adjusted to the center of gravity, and an imaging error is corrected by using an Equally Sloped Tomography (EST) algorithm, thereby achieving a resolution of 2.4 ° and enabling observation at the atomic level.

  In the observation and analysis methods that can be performed at the atomic level, attempts to improve the resolution by correcting aberrations are constantly being made. For example, Patent Document 1 can be cited.

  Further, as a technique corresponding to the above (2), there is an electron diffraction pattern observation.

JP 2010-153320 A

Sandra Van Aert1, Kees J. et al. Battenburg, Martin D. See Rossell, Rolf Erni, and Gustaaf Van Tenderoo. Three-dimensional atomic imaging of crystalline nanoparticles, nature09741, 2011 M. C. Scott, Chien-Chun Chen, Matthew Mecklenburg, Chun Zhu, Rui Xu, Peter Ercius, Ulrich Dahmen, B.S. C. Regan, and Jianwei Miao. Electron tomography at 2.4 angstrom resolution, nature 10934, 2011.

  However, the methods described in Non-patent Document 1 and Non-patent Document 2 each obtain an image using a STEM, and therefore can only specify the position of each atom constituting a nanoparticle. It does not reflect the finite periodicity in the area of the meter order.

  Therefore, in the analysis methods of Non-Patent Documents 1 and 2, the material properties (physical properties) based on the wave number space, that is, the finite periodicity and the obtained image are associated with each other, The size and shape, as well as the physical properties, chemical properties and chemical reactions of the nanoparticles cannot be fully discussed.

  That is, since the interpretation based on the wave number space cannot be performed, it is difficult to analyze the momentum space based on the canonical conjugate based on the quantum theory.

  Therefore, the above prior art does not enable analysis of dramatic changes in electrical conductivity caused by regular inflection or bending of the band structure caused by finite periodicity, radicalization of chemical reactions, etc. Analysis of finite periodicity does not directly contribute to the development of material science.

The electron diffraction pattern observation of the above (2) can obtain a diffraction pattern of particles on the order of micrometers, but conventionally, it is assumed that the measurement sample has infinite periodicity, and the Laue electron diffraction Atomic positions were specified by the main spot, ignoring the sub-spot.
Therefore, it is not intended to analyze a finite periodicity as in the method (1).

  As described above, clustering and amorphization have progressed, and the analysis of nanoparticles in which the finite periodicity of the atomic arrangement is about to collapse, and the physical properties and chemical properties based on the analysis results of the above nanoparticles, and those physical properties and Analysis based on the time evolution of chemical properties cannot be said to be achievable even with the latest science and technology.

The present invention has been made in view of the problems and future prospects of such conventional technology, and has a finite periodicity, enabling analysis of fine particles such as so-called nanoparticles having finite periodicity, The purpose is to contribute to the analysis of various fine particles such as nanoparticles that will be developed in the future.
It is another object of the present invention to enable analysis of fine particles in which clustering and amorphousization have progressed and even finite periodicity has begun to collapse.

  Furthermore, the progress of analytical technology in the 21st century largely depends on the utilization of quantum beams, ie, synchrotron radiation, electron beams, and neutron beams, and it is natural that recent atomic-level structural analysis is performed using synchrotron radiation. Accordingly, it is an object of the present invention to enable analysis and analysis in cooperation with various quantum beam analysis results including the above-mentioned synchrotron radiation.

  In other words, the present invention predicts and analyzes the shape of fine particles such as so-called nanoparticles having finite periodicity, and the shape of fine particles having finite periodicity that is about to collapse, at the atomic level resolution for each plane orientation. Microanalysis method of fine particles capable of shape analysis capable of solving even physical properties, chemical properties and chemical reactions of the fine particles as well as the fine particles, an analyzer used for the microscopic analysis method, and An object of the present invention is to provide a method of operating a microscope of an analyzer used for the method of analyzing a microscope.

  The present inventors have conducted intensive studies to achieve the above object, and as a result, focused on the essential properties of the Laue function, and ignored when analyzing a material that can be regarded as having infinite periodicity. The inventor has found that the use of the Laue electron diffraction subspot can realize microscopic observation in consideration of the finite periodicity of fine particles such as nanoparticles, and have completed the present invention.

That is, the present invention is based on the above findings, the microscopic analysis method of the fine particles of the present invention,
(1) a beam irradiation step of irradiating fine particles with a parallel beam;
A detection step of imaging a plurality of electron diffraction patterns by changing the imaging angle and / or the size of the Ewald sphere for the fine particles;
An array confirmation step of analyzing the diffraction spot of the electron diffraction pattern to determine the shape of the fine particles,
Including
The arrangement confirming step compares an input model predicting the size and shape of the fine particles with an array of Laue sub spots of the plurality of electron diffraction patterns, and determines the shape of the fine particles by correcting the input model. It is characterized by that.
(2) The fine particles have a cluster portion and / or an amorphous portion .
(3) In the arrangement confirmation step, the input model is compared with the arrangement of the main spots and Laue sub-spots of the plurality of electron diffraction patterns .
(4) The input model is constructed using a lattice constant of a finite periodic part, a space group, and partial atomic coordinates .
(5) The sequence confirmation step includes an atomic position determination process,
The atomic position determination processing is characterized in that the position of an atom is determined using a Fourier transform and an inverse Fourier transform.
(6) The arrangement confirmation step includes an analysis of the intensity of the Laue sub spot.
(7) An electron diffraction pattern is photographed by a parallel electron beam system after confirming a photographing angle of fine particles by a convergent electron diffraction system.

Further, the microscopic analyzer for fine particles of the present invention,
( 8 ) An electron beam irradiation unit for irradiating the sample with a parallel electron beam, an intermediate lens, a projection lens, a zero-order reflection blocking mechanism for blocking zero-order reflection, a main spot blocking mechanism, and a Laue sub-spot of the sample. And a detector for detecting
The electron beam irradiator has a converging lens for converging an electron beam, and the converging lens can form an opening angle of less than 1.0 mrad.

According to the present invention, since the Laue sub-spot analysis between the main spots is performed to confirm the shape of the fine particles, the atomic arrangement cannot be regarded as having an infinite periodicity, that is, nanoparticles having a so-called finite periodicity It is possible to predict and analyze the shape of fine particles such as at the atomic level resolution for each plane orientation, and to clarify physical properties, chemical properties and chemical reactions by the shape analysis.
Furthermore, even within a finite region on the order of nanometers, it becomes possible to predict and analyze the shape of fine particles such as nanoparticles whose periodicity is about to be destroyed at an atomic level resolution.

FIG. 3 is a graph illustrating a Laue function according to the first embodiment of the present invention. FIG. 1 is a schematic diagram illustrating a nanobeam electron beam optical system according to a first embodiment of the present invention. It is the graph which showed the Laue spot concerning a 1st embodiment of the present invention typically, and the projection figure. It is a flow chart figure of Laue spot arrangement analysis concerning a 2nd embodiment of the present invention. It is a flow chart figure of Laue spot arrangement analysis concerning other 2nd embodiments of the present invention. FIG. 11 is a schematic diagram illustrating a nanoparticle shape and an electron diffraction pattern according to a third embodiment of the present invention. FIG. 14 is a schematic diagram illustrating a nanoparticle shape and an electron diffraction pattern according to a fourth embodiment of the present invention. It is the schematic diagram showing the flowchart and nanoparticle shape which concern on 5th Embodiment of this invention. It is a graph figure concerning a 6th embodiment of the present invention. It is a schematic diagram and a graph diagram concerning a seventh embodiment of the present invention. It is the schematic diagram and graph diagram which concern on other 7th Embodiment of this invention. It is a mimetic diagram concerning an 8th embodiment of the present invention. It is a photograph figure showing an example. It is a photograph figure which shows a comparative example. It is a photograph figure and a spectrum figure showing an example. It is a photograph figure and a spectrum figure (an enlarged view) showing an example.

<Microscopic analysis of fine particles>
The microscopic analysis method for fine particles of the present invention will be described, but the present invention is not limited to the following embodiments. The present invention is characterized by focusing on a weak Laue subspot that has been ignored or excluded in the past.

[principle]
The diffraction intensity I of fine particles such as nanoparticles having periodicity in a region on the order of nanometers, so-called nanoparticles having finite periodicity, is considered excluding the influence of crystal structure factors (space group) and thermal vibration of atoms. And the square of the magnitude of the atomic scattering factor f and the square of the magnitude of the Laue function, as shown in (Equation 1).

[Equation 1]
^{_{I α | f | 2 · |}} L a (q a) | 2 · | L b (q b) | 2 · | L c q c) | 2 ( Equation 1)
_{L n (q n) = |} (sinπN n q n d hkl) / (sinπq n d hkl) | 2
n: a, b, or c lattice axis
N _n : number of lattices in n lattice axis direction q _n : scattering vector in n direction h, k,?: Index d _hkl : lattice plane spacing

The atomic scattering factor is a factor depending on the atom, and does not depend on whether it is infinite periodicity or finite periodicity. Therefore, it is sufficient to focus on the Laue function in order to discuss the diffraction intensity when there is a finite periodicity.
There are various factors to be considered in actual experiments, such as the Debye-Waller factor. Here, in order to simplify the discussion and clarify the point of interest of the present invention, the above-mentioned Debye-Waller factor will not be referred to.

FIG. 1 is a graph showing the diffraction intensity I in quantum beam diffraction. As an example, as a simple cubic lattice, the lattice spacing d ₁₀₀ 3 Å (hence, reciprocal lattice that is in the case of 1/3 Å ^-1) shows the diffraction intensity.

Assuming that the fine particles to be measured have infinite periodicity, as shown in FIG. 1A, since the lattice vector of the lattice is 1 / 3Å- ¹ , the scattering vector is 1/3 and its integer. Only the main spot 101 satisfying the double appears.
Thus, assuming that the fine particles to be measured have infinite periodicity, no diffraction spot (peak) appears between the main spots, and the spot intensity becomes zero.

  The diffraction phenomenon is a phenomenon that spreads three-dimensionally, but the diffraction pattern is a cross section of the diffraction pattern and has two-dimensional data, and thus is usually called a spot. In FIG. 1A, since the two-dimensional data is converted into a graph (that is, one-dimensional data), the main spot 101 in FIG. 1A should be described as a main peak.

  Next, consider the case of so-called fine particles having finite periodicity. As an example, consider a case where six unit cells are arranged in the a-axis direction. Note that a similar discussion can be made in the b-axis and c-axis directions.

  FIG. 1B shows the diffraction intensity in the case where the fine particles to be measured are nanoparticles having a size of several nanometers, each of which has a length of 3 ° in each axial direction of the unit cell and is constituted by six unit cells. I is shown.

As shown in FIG. 1B, a plurality of Laue sub-spots 102 appear between the diffraction main spots 101 of the wave vector zero [Å ⁻¹ ] and ３ [Å ⁻¹ ].
The length of each of these nanoparticles in the axial direction is 3Å × 6 = 18Å, which is about 2 nm. The number of d ₁₀₀ lattice plane is seven, inclusive, the number of lattice planes other than the end faces are five.

FIG. 1C shows an enlarged view of the Laue subspot surrounded by a frame line in FIG. As shown in FIG. 1 (c), Laue sub-spots in the form corresponding to the number of d ₁₀₀ lattice planes of the internal nanoparticles (except both end portions), five Laue sub-spot 102 is generated.
The peaks of the Laue sub-spots 102 are arranged at equal intervals according to the phase difference of the waves diffracted on each surface, while the spot intensity of each Laue sub-spot 102 is different.

The number of appearances of the Laue sub-spots, the distance between the Laue sub-spots, and the intensity of the Laue sub-spot depend on the number of lattices in the plane direction.
Since the orientation of the lattice differs for each index h, k,?, That is, for each scattering (wave number) vector, it is possible to know the length from the center to the end in each of the various directions of the nanoparticle of interest. For this reason, it is possible to obtain not only the mere size of the nanoparticles but also a detailed shape from the Laue subspot.

  In addition, since the wave number vector of the exposed surface is known, the wave number vector and the momentum as the canonical conjugate amount are simultaneously known. This enables analysis / analysis in cooperation with other analyzers, and also enables detailed knowledge of the material properties (physical properties) of the nanoparticles for each direction.

[Microanalyzer]
Next, a microanalyzer used in the microanalysis method of the present invention will be described. The present invention is to detect and analyze Laue sub-spots by a Laue sub-spot detection optical system that detects a Laue sub-spot which is a weak diffraction spot from nanoparticles.
The Laue subspot detection optical system detects an Laue subspot by collimating an electron beam and steadily irradiating the nanoparticles.
FIG. 2 is a schematic diagram showing a Laue sub-spot detection optical system of the microscopic analyzer according to the present invention.

The measurement / analysis system in the microanalyzer according to the present invention is based on a conventional electron microscope, and includes an electron beam irradiator for irradiating a sample with a parallel electron beam, an intermediate lens, a projection lens, and a zero-order reflection. It has a zero-order reflection blocking mechanism for blocking light, and a detector for detecting a Laue sub-spot of the sample, and has a main spot blocking mechanism for blocking a main spot as required, an objective diaphragm, a field limiting diaphragm, and the like.
The electron beam irradiation unit has a converging lens for narrowing the electron beam, and the converging lens forms an electron beam with high parallelism. It is desirable that the opening angle 2α is as small as possible, but it cannot be set to 0 °.

  The electron beam irradiation section forms a parallel electron beam and irradiates the sample with the parallel electron beam. It is preferable that the parallel electron beam irradiating the sample is applied to only the particles to be measured without excess or deficiency so that the transmission electron beam is not generated from a part other than the nanoparticles.

  Therefore, it is important to narrow down the diameter of the parallel electron beam to a diameter (nanoparticle diameter to nanoparticle diameter × 1.1) according to the size of the nanoparticle to be measured, and to increase the parallelism of the electron beam. The electron beam is narrowed down by a converging lens having an aperture angle of about 0.5 to 1.0 mrad of the parallel electron beam indicated by 2α in FIG.

  As the convergent lens 112, it is preferable to use a convergent lens that can form an electric field and a magnetic field that can be narrowed down while maintaining parallelism, instead of a normal convergent lens. To ensure parallelism and convergence, the converging lens 112 may have a multi-stage configuration.

  The narrowing of the parallel electron beam can be performed by adjusting the electromagnetic field of the converging lens 112, and the narrowing can be performed while maintaining the parallelism of the electron beam.

  The parallel electron beam applied to the sample passes through the sample, passes through the first intermediate lens 113, the second intermediate lens 114, and the projection lens 115, reaches the projection surface 116, and captures an electron diffraction pattern.

  Since the present invention detects Laue sub-spots, it includes a zero-order reflection blocking mechanism 117 so that an electron beam does not enter the detector 119. By excluding the zero-order reflection, the saturation of the detector is suppressed, and the S / N ratio of the detector can be used only for detecting the main spot and Laue sub-spot.

  In order to determine the shape and size of the nanoparticles, it is most important to detect a weak Laue subspot. In particular, when it is assumed that the nanoparticles are quite large and a large number of Laue sub-spots are detected, or that the number of Laue sub-spots is weak, the electron beam is prevented from entering the detector 119 according to the dynamic range of the detector 119. It is preferable to insert the main spot cutoff mechanism 118 into the main spot and exclude the main spot.

  In other words, when the nanoparticles to be measured have a diameter of several nanometers or even more than tens of nanometers, the number of Laue sub-spots increases from several tens to sometimes hundreds or more, and the Laue sub-spots increase. The strength becomes extremely weak.

  The main spot cutoff mechanism 118 may also cover the Laue sub-spot that has appeared near the main spot. However, when it is necessary to exclude the main spot, it is a case where the nanoparticles to be measured are large, so that it is possible to detect a large number of weak Laue sub-spots that appear near the middle between the main spots.

  Therefore, by using the main spot blocking mechanism 118, even if the Laue sub-spots are weak, the shape and size of the nanoparticles can be detected by detecting a large number of weak Laue sub-spots that appear near the middle between the main spots. You can decide.

  When it is desired to know the intensity of the diffraction spot accurately, the intensity can be measured without using the main spot blocking mechanism 118.

  The detector 119 for detecting a Laue sub-spot from a sample preferably has a high S / N ratio in order to capture a weak Laue sub-spot. Specifically, when the S / N ratio is 3 digits or more, Preferably, there is.

  As the detector, any device capable of taking a photograph may be used, and examples thereof include devices such as a film, a CCD (Charge Couple Device), an Imaging Plate, and a Pixel Apparatus for the SLS (Pilatus (trade name)).

The film, CCD, Imaging Plate, and the S / N ratio of Pilatus ^are each a 10 ^2, 10 5, 10 about ^6, from the viewpoint of S / N ratio, the detector in the order of CCD <Imaging Plate <Pilatus It can be said that it is suitable.

The microscope analyzer of the present invention can use a selected area stop.
The convergence and the parallelism of the electron beam are in conflicting relations. If priority is given to the parallelism, the diameter of the electron beam increases and it becomes difficult to apply the electron beam only to the nanoparticles.
In order to create a state where only the nanoparticles are irradiated with the electron beam while maintaining the parallelism of the electron beam, it is preferable to insert the selected area stop 120 between the sample and the intermediate lens.

  When the selected area stop is a selected area stop composed of fine holes prepared for nanoparticles, it is possible to create a state in which an electron beam is applied only to nanoparticles while maintaining parallelism.

  Hereinafter, the microscopic analysis method for fine particles of the present invention will be described in detail with reference to embodiments, but the present invention is not limited to the following embodiments.

  The microscopic analysis method for fine particles of the present invention includes a beam irradiation step of irradiating a sample with a parallel beam, a detection step of detecting a diffraction spot from the sample, and an array confirmation of analyzing the diffraction spot to determine the shape of the fine particle. The sequence confirmation step includes a Laue subspot analysis between main spots to determine the shape of the fine particles.

(First embodiment)
The feature of the first embodiment is that a large number of electron diffraction patterns from low-order reflection to high-order reflection are taken and used complementarily to each other, so that the precise size and size of the nanoparticles can be improved. It is in the sequence confirmation step of analyzing and determining the shape.

The sequence confirmation step will be described.
The arrangement confirmation step is a step of calculating the interval of the Laue subspot arrangement from the electron diffraction pattern in which the Laue subspot is photographed, and determining the shape of the particles.

FIG. 3 is a schematic diagram showing Laue sub-spots. FIG. 3A is a graph showing an electron diffraction pattern which is two-dimensional data in a three-dimensional graph, and FIG. 3B schematically shows an electron diffraction pattern obtained by observation with a transmission electron microscope. It is a virtual photograph figure.
FIG. 3 (b) schematically shows the diffraction spot intensity in a two-dimensional manner reflecting the spot diameter, and a circle in FIG. 3 (b) indicates the spot diameter (diffuse scattering). Note that this is not a representation.

  Since the actual nanoparticles have a length of several nanometers in each direction, they should be composed of several tens of lattice planes, and the number of Laue sub-spots existing between the main spots is several tens. It should be an individual. However, since it is difficult to show several tens of sub-spots, the number of lattice planes is shown in FIG.

The first embodiment utilizes a real lattice-reciprocal lattice transformation.
In this embodiment, the incident direction of the electron beam is arbitrary. Since the maximum detection accuracy of the length is equal to the lattice spacing, the electron diffraction pattern in which the lattice spacing is as small as possible within the range where the Laue sub-spot can be detected, that is, the scattering vector spot with the largest possible index is shown. Selection is preferred.

A simple cubic lattice (Pm-3m) having a lattice constant of 3 ° will be described as an example.
As an example, consider a Laue sub-spot when the incident direction [1−32] of an electron beam is accidentally captured by an electron microscope.

Although it is originally correct that the azimuth (scattering vector) of the electron beam is described without parentheses, in order to make the distinction between the azimuth and the plane index clear and to ensure the legibility, in this specification, the brackets are used. In addition, the azimuth and the scattering vector will be represented.
In addition, when indicating a general term of an azimuth, the azimuth is expressed by angle brackets in accordance with a custom in the field of electron microscopes.

As a scattering vector perpendicular to [1-32], for example, there is [2 2-3].
Further, [5 7 8] can be mentioned as a scattering vector perpendicular to both [1-32] and [2 2-3].
In FIG. 3A, the [2 2 -3] direction is referred to as a scattering vector (horizontal direction), and [5 7 8] is referred to as a scattering vector (vertical direction).

The lattice spacing of [22-3] is 0.802 °, that is, about 0.8 °. In the horizontal scattering vector, five Laue sub-spots 102 are visible between the main spots 101.
Therefore, it can be determined that five lattices are arranged in the [2 2 -3] direction and that 5 + 1 = 6 lattice planes exist. From this result, it is determined that the distance of the nanoparticles from the center to the [2 2 -3] lattice plane is 0.8 × 6 = 4.8 °.

  Actual nanoparticles generally belong to various space groups and include many types of atoms. However, if the nanoparticles in the present embodiment are a simple cubic lattice composed of a single atom, For example, since the number of planes and the number of atoms are the same, the number of atoms in the [2 2 -3] direction is determined to be 6.

  Similarly, in the scattering vector in the vertical direction, six Laue sub-spots 102 are visible between the main spots 101. That is, it can be determined that [5 7 8] is made up of 6 + 1 = 7 planes.

  The lattice spacing of [5 7 8] is 0.255 °, that is, about 0.26 °. The length of the nanoparticles in the [5 7 8] direction (vertical direction) is 0.26 × 7 = 1.8 °, and assuming that the nanoparticles are a simple cubic lattice composed of a single atom, the vertical direction Means that seven atoms are aligned.

By analyzing the Laue sub-spots in this way, the length of the nanoparticles in two directions can usually be known from one electron diffraction pattern.
Therefore, by photographing n electron diffraction patterns, it is usually possible to know the length in the 2n direction and the number of atoms in that direction.

Also, considering the space group of the periodic part and the partial coordinates of each atom, the position of each atom in that direction of the nanoparticle can be determined with the lattice spacing being the maximum positional accuracy.
Therefore, by photographing about 100 electron diffraction patterns, it is possible to know the length of the nanoparticles and the partial atomic coordinates of each atomic species in almost all directions (all regions) of the nanoparticles, but the analysis is performed. The number of electron diffraction patterns may be appropriately selected according to the required accuracy.

(Second embodiment)
Next, a second embodiment of the microscopic analysis method for fine particles according to the present invention will be described.
In the first embodiment, it has been shown that by analyzing the Laue subspot of the electron diffraction pattern, the length of the nanoparticle and the partial coordinates for each atomic species can be determined.

  In the method of the first embodiment, the relative positional relationship between the atoms and the partial atomic coordinates of the periodic portion are known, so that it can be said that the accuracy of determining the atomic position is high. However, since the accuracy with respect to the length is on the order of the lattice spacing, it is not always sufficient.

  In the present invention, since the Laue spot having a weak strength is used, the spot shape is likely to be diffused. In the arrangement confirmation step of the present embodiment, in order to perform the real lattice-reciprocal lattice transformation even in the state where the spots are diffused, the arrangement is further confirmed using the Fourier transformation-inverse Fourier transformation, and the atom position determination step of determining the position of the atom is performed. Is added.

  FIG. 4 is a flowchart of Laue spot array analysis according to the second embodiment of the present invention. FIG. 4 shows an example of the form of the procedure when the partial atomic coordinates can be simply determined.

  In the second embodiment, an input model S101 is constructed, a Fourier transform S102 is applied to the input model S101, and an array of main spots and Laue sub-spots of the input model is calculated.

  The input model predicts the size and shape of the nanoparticle to be measured.Specifically, the size of the nanoparticle and the approximate shape of the nanoparticle, such as whether the shape is close to a cube or close to a rectangular parallelepiped, You should have an idea.

  In constructing the input model, the lattice constant, space group, and partial atomic coordinates of the finite periodic part of the nanoparticles are used as initial data. Although it is not always necessary, it is preferable to construct an input model 101 by taking a real space image of the nanoparticles in advance and comprehensively considering the size and shape of the nanoparticles.

  Then, the Laue sub-spot arrangement confirmation S103 is performed, and the arrangement of the main spot and the Laue sub-spot of the input model is compared with the arrangement of the main spot and the Laue sub-spot of the actually captured electron diffraction pattern to correct the input model. To get closer to the actual fine particles.

  Specifically, the Laue subspots of the electron diffraction pattern and the Laue subpot obtained by applying the Fourier transform S102 to the input model S101 have the same interval and number of Laue subspots. , The lattice spacing of the input model, and the number of the lattice planes.

  Normally, there are two directions of scattering vectors perpendicular to the incident azimuth per electron diffraction pattern. Therefore, the lattice plane spacing and the number of lattice planes are determined for each scattering vector direction. The length of the scattering vector can be known from the lattice plane interval and the number of lattice plane intervals.

  A similar operation is performed on the second electron diffraction pattern. Since the azimuth of the second electron diffraction pattern is different from that of the first electron diffraction pattern, the lattice spacing and the number of lattice spacings can be obtained for scattering vectors different from the first one.

  By applying the same operation to all the obtained electron diffraction patterns, the lattice spacing, the number of lattice spacings, and the length in almost all directions of the nanoparticles are determined for each direction. The spot arrangement confirmation S103 ends.

  Next, by applying the inverse Fourier transform S104 to the Laue subspot arrangement of the confirmed input model, the arrangement of atoms is restored from the lattice spacing and the number of lattice planes, and model output S105 is performed.

  A comparison S106 between the input model and the arrangement of atoms (size and shape of nanoparticles) of the output model output in the model output S105 is performed to obtain a new input model S101 having a more appropriate size and shape. To construct.

  Such a Fourier transform S102 → Laue spot arrangement confirmation S103 → Fourier inverse transform S104 → model output S105 is repeatedly repeated until the structure between the input model and the output model is approximately the same, and the model output S105 I do.

  As described above, if the Fourier transform-inverse Fourier transform is applied and the Laue sub real lattice-reciprocal lattice transform is repeatedly performed, the arrangement interval of the Laue sub spot array can be accurately obtained even when the spots are diffused. .

  That is, the shape of the input model and the shape of the output model can be directly compared by obtaining the real lattice image by performing the inverse Fourier transform. That is, by observing and comparing images (images), it can be confirmed from the viewpoint of human eyes. In addition, the evaluation can be made not by the abstract concept of the Laue subspot arrangement interval in the reciprocal lattice (inverse space) but by the real concept of the coordinates of each atom of a real particle. As a result, the position of each atom in the nanoparticle can be determined more accurately than in the first embodiment, and its validity can be determined and evaluated.

  Next, an example of an embodiment in which the number of steps is increased and the procedure becomes more complicated, but more suitable for actual use of an electron microscope and in which partial atomic coordinates can be determined with high accuracy will be described with reference to FIG.

  A plurality (n) of electron diffraction images of a simple plane index such as the (100) plane are taken in advance. From the obtained electron diffraction pattern, the camera length is corrected to match the crystal structure data used in the input model construction S201, and lattice constant correction S202 is performed.

  Thereafter, the Fourier transform S203 is performed on all the n electron diffraction patterns, and the Laue spot arrangement confirmation S204 similar to the case of FIG. 4 is performed. By applying the inverse Fourier transform S205 to all data after the Laue sub spot arrangement confirmation S204, partial coordinates of each atom are calculated for each lattice plane direction.

Next, an atomic position determination process S206 is performed.
The above-described atomic position determination processing S206 is processing for determining partial atomic coordinates of each atom.
First, a certain lattice plane (which will be referred to as a first lattice plane) is selected, has a different plane orientation from the first lattice plane, and has the same partial coordinates as the first lattice plane. A second lattice plane containing a power atom (to be referred to as a target atom) is selected.

  The partial coordinates of the atom of interest on the second lattice plane are shifted in the direction of the scattering vector such that the partial coordinates of the atom of interest match based on the partial coordinates of the first lattice plane.

  By determining the partial coordinates of the target atom in this manner, the target atom and the portion of the atom that should maintain the positional relationship with the target atom with the same accuracy as the crystal structure data used in forming the input model. Coordinates can be determined.

  A similar operation is performed on n electron diffraction patterns that have been photographed. Although there are many electron diffraction patterns, depending on the number of photographed pictures, there are a plurality of ways to select two types of lattice planes.

  Then, the partial coordinates of the atom of interest may be different for each combination of the selected lattice planes due to the influence of errors in the measurement system and the like. Similarly, the partial coordinates of the atoms that should maintain the positional relationship with the target atom may differ depending on the selected combination of lattice planes.

  Therefore, in the atomic position determination processing, a series of processing from the Fourier transform S203 to the input model and comparison S208 is repeated so as to minimize this error, and an optimized structural model is obtained. This makes it possible to determine the partial atomic coordinates of each atom with high accuracy even when there is a mechanical error in the measurement system.

FIG. 5B shows an image of the nanoparticles determined in this manner.
The length of each side of the (hkl) plane depends on the scattering vector (k ^* and l ^{* for the} a-axis direction, h ^* and l ^{* for} the b-axis direction, and h ^* and k ^{* for the} c-axis direction). Decided. Therefore, in the case of high-order reflection, h ≠ k ≠ l, so that each surface is represented by a scalene triangle.

As described above, the partial coordinates of the target atom on the second lattice plane are shifted in the direction of the scattering vector so that the partial coordinates of the target atom coincide with each other with reference to the partial coordinates of the first lattice plane. The coordinates are determined.
Note that atoms do not necessarily exist at the points where the lattice planes intersect, but in FIG. 5B, the atoms are drawn at the points of intersection (vertices of a triangle) for convenience.

  In the analysis method described above, the process of performing optimization after selecting two types of lattice planes has been described. However, it is also possible to select n lattice planes and perform optimization at once. By performing the optimization by selecting one of the optimization methods, it becomes possible to determine the coordinates of each atom in the nanoparticle with higher accuracy than in the first embodiment.

As a method of capturing an electron diffraction pattern at a stretch, there is an electron backscatter diffraction (EBSD).
EBSD is a method of capturing a Kikuchi line by setting the optical system to a convergent electron diffraction system in a transmission electron microscope. The use of the Kikuchi map showing the Kikuchi line enables the orientation of nanoparticles to various directions with high accuracy. Since this technique has already been put to practical use in a scanning electron microscope, it will not be described in detail.

  It is quite difficult to apply the above EBSD to a transmission electron microscope. In particular, it is substantially impossible to apply the present invention to a Laue subspot detection optical system which is a parallel electron beam system of the present invention for detecting a Laue subspot. Therefore, after photographing with the Laue sub spot detection optical system, the optical system is temporarily changed to a convergent electron diffraction system.

  That is, the angle α is temporarily widened by changing the distance between the converging lens 112 and the sample 111 in FIG. This operation makes it possible to track the Kikuchi line found in the diffraction ring.

  Tracking of the Kikuchi line can be automated, as is already done with scanning electron microscopes. By tracking the Kikuchi line, the sample is automatically tilted in various directions with high accuracy. When the sample position is determined, the optical system is returned from the convergent electron diffraction system to the Laue sub-spot detection optical system which is a parallel electron beam system.

  By automatically performing the conversion between the Laue subspot detection optical system, which is a parallel electron beam system, and the conventional convergent electron diffraction system many times, the shape of the nanoparticles and the positions of the constituent atoms can be determined with high accuracy. It can be obtained automatically.

(Third embodiment)
Next, a third embodiment of the microscopic analysis method for fine particles according to the present invention will be described. In the third embodiment, the shape of nanoparticles is determined by two electron diffraction patterns.

  Nanoparticles are three-dimensional objects, while electron diffraction patterns are two-dimensional data. Therefore, it is difficult to determine the shape of the nanoparticles from one electron diffraction pattern.

  However, if there are two electron diffraction patterns, 2 (dimension) +2 (dimension) information can be obtained and the amount of information exceeds 3 (dimension) of the object. Then, the size and shape of the nanoparticles can be discussed.

FIG. 6 is a schematic diagram illustrating a nanoparticle shape and an electron diffraction pattern according to the third embodiment of the present invention. 6B and 6C, the main spots are indicated by hatched circles, and the Laue sub spots are indicated by black circles.
Note that the radius of the Laue sub-spot schematically represents the spot intensity, and does not mean that the spot has spread due to influence such as distraction.

  Analysis of the nanoparticles in the input model assuming that the finite periodic part of the nanoparticles is formed by a simple cubic lattice and consists of three types of lattice planes, {100}, {110}, and {111} An example will be described.

  As for the {100} and {110} planes, the plane shape cannot be defined because the length of the corresponding side becomes 1/0, that is, infinity, such as the {100} and {110} planes.

  Therefore, it is assumed that the {100} plane is square because it includes two 0s and two sides are indeterminate. Since the {110} plane includes one 0 and one side is indefinite, it is assumed that one side has a degree of freedom and is a rectangle. The shape of the {111} plane is an equilateral triangle because the exponent does not include 0 and the exponent values are all 1 and equal to each other.

  It is assumed that the size (area) of each lattice plane of the input model is {100}> {110}> {111} as shown in FIG. That is, square> rectangle> regular triangle. The length for each orientation of the nanoparticles is inversely related to the area.

Therefore, an input model having the longest <111> orientation of the nanoparticle, followed by the shortest <110> and <100> orientations was selected.
From another viewpoint, the input model is set when the number of {111} planes arranged in the <111> direction is the largest, and then the {110} plane and the {100} plane are largest.

  The number of Laue sub-spots for each direction is determined by the number of lattice planes. As shown in FIGS. 6B and 6C, the number of sub-spots in each direction in the present embodiment is in the order of <111>, <110>, and <100>.

  The spot intensities in FIG. 6B and FIG. 6C cannot be determined unless the crystal structure of the finite periodic portion is specified, and are therefore convenient for explaining the present embodiment. The diameter of the nanoparticles is usually several nm. In this case, the number of lattice planes, that is, the number of Laue sub-spots, is from tens to several tens. Since it is difficult to show such a large number of spots, it is different from the actual one, but is illustrated with several Laue sub-spots.

When an electron beam is incident in the [001] direction as shown in FIG. 6A, an electron diffraction pattern including {100} and {110} reflections as shown in FIG. 6B can be obtained.
Similarly, if the electron beam incident direction is the [11-2] direction, an electron diffraction pattern including {110} and {111} reflections can be obtained.

For each of {100}, {110}, and {111}, the number of lattice planes of the finite periodic part can be known from the number of Laue sub-spots.
Therefore, the shape of the assumed nanoparticle composed of three types of {100}, {110}, and {111} is determined. Further, {110} is included in both electron diffraction patterns.
Then, by applying the analysis method described in the second embodiment, it is possible to determine the partial coordinates for each direction.

  As described above, according to the present embodiment, it is possible to obtain the shape of nanoparticles having finite periodicity and the partial coordinates of each atom forming the shape even when only two electron diffraction patterns can be captured. Can be.

(Fourth embodiment)
Next, a fourth embodiment of the microscopic analysis method for fine particles according to the present invention will be described.
The fourth embodiment enables a three-dimensional shape to be discussed even when only one electron diffraction pattern is obtained.

FIG. 7 is a schematic diagram illustrating a nanoparticle shape and an electron diffraction pattern of an input model according to the fourth embodiment of the present invention.
Assuming the incidence of [001], information on the (001) and (00-1) planes cannot be obtained, so the shape of this plane is represented by an octagon. The {100} planes other than the (001) and (00-1) planes were represented by squares. The {110} plane was rectangular.

Assuming that the lengths of the nanoparticles in the [100] and [110] directions are substantially equal, the ratio of the number of lattice planes is √ (1 ² + ¹² + 0) = √2. The number of Laue sub-spots for <100> was set to 2 and the number of Laue sub-spots for <110> orientation was set to 3 so that the number of Laue sub-spots was close to the ratio of the number of lattice planes.
Note that this is only for convenience in drawing, and is not the actual number of Laue sub-spots and the actual intensity.

For the [001] direction, the length of the nanoparticles in this direction is estimated using the intensity of the electron beam. The estimation method is as follows: 1) Ratio between direct beam (0th order reflection spot) and main spot (ratio between transmitted ray and scattered ray)
2) Ratio of main spot to Laue spot 3) Calculation from EELS.

  As described above, according to the present embodiment, it is possible to discuss the shape of nanoparticles having finite periodicity from one electron diffraction pattern.

(Fifth embodiment)
Next, a fifth embodiment of the microscopic analysis method for fine particles according to the present invention will be described.
The fifth embodiment can be applied to the case where a large number of electron diffraction patterns are obtained as in the first and second embodiments, but the obtained electron diffraction pattern is different from that in the third embodiment. The present invention relates to an analysis method that enables a three-dimensional shape to be discussed even when only two electron diffraction patterns are obtained, as in the fourth embodiment.

FIG. 8 is a schematic diagram illustrating a flowchart and a nanoparticle shape according to a fifth embodiment of the present invention. A Laue spot arrangement in two directions can be obtained from one electron diffraction pattern.
Therefore, although it is possible to discuss the shape with two types of lattice planes, the present embodiment shows an analysis example in the case where only one type of lattice plane is used.

For the sake of simplicity, it is assumed that the finite periodic part of the nanoparticle consists of a simple cubic lattice.
As shown in FIG. 8A, first, an input model construction S301 is constructed. Since it is an input model, as shown in FIG. 8B, the shape of the nanoparticle is the same in any of the directions of the a, b, and c axes, that is, it is a cube. A simulation is performed in which the lattice constant and the number of lattice planes of the cubic nanoparticles are variously changed, and an electron diffraction pattern at the lattice constant and the number of lattice planes is reproduced.

  This electron diffraction pattern is compared with an electron diffraction pattern actually obtained by observation with a transmission electron microscope. Then, as shown by a broken line in FIG. 8B, for example, the result is obtained that the number of lattice planes in the [100] direction is larger than that in the [010] and [001] directions. . In this case, the nanoparticles are not cubes but cuboids extending in the [100] direction.

  Next, the above-described simulation is started from an input model constituted by {110} as shown in FIG. 8C, for example, and the number of lattice planes for each direction is determined. The length of the nanoparticle in that direction is determined according to the number of lattice planes.

  Similarly, as shown in FIG. 8D, a simulation may be performed using nanoparticles composed of {111} as an input model. Of the {100}, {110}, and {111} models, the one that can best explain the electron diffraction pattern obtained in the experiment is selected.

Since this is an input model, the calculation starts with an input model calculation of nanoparticles composed of one kind of lattice plane.
As a result, for example, it is determined that it is more appropriate to consider that the nanoparticles are composed of the {100} and {110} planes. Even if only one electron diffraction pattern was obtained in the experiment, the electron diffraction pattern should have obtained electron diffraction spots for two directions, that is, two types of lattice planes.

  Assuming that the lattice planes are {100} and {110}, an input model of the nanoparticles composed of the {100} and {110} planes is constructed, and a new calculation (repetition) is performed as shown in FIG. ) You just have to do the calculations.

  When the shape of the nanoparticles is determined by such a simulation, it is clear what the main exposed surface is. As is well known, various chemical properties such as catalytic activity are significantly different for each exposed surface of the nanoparticles. Since the preferred growth surface of the nanoparticles can be specified, the present embodiment is expected to exert its power in developing a material having desired characteristics.

(Sixth embodiment)
Next, a seventh embodiment of the microscopic analysis method for fine particles according to the present invention will be described. The seventh embodiment is a preferred embodiment when a large number of electron diffraction patterns are obtained as in the first and second embodiments.

FIG. 9 is a schematic diagram illustrating a flowchart and a nanoparticle shape according to a seventh embodiment of the present invention.
In this embodiment, as shown in FIG. 9A, electron diffraction patterns of nanoparticles composed of 2 to n types of lattice planes are used by using data of many lattice planes obtained from a large number of electron diffraction patterns. The simulation S402 and the comparison S403 with the electron diffraction pattern obtained by the experiment are repeatedly subjected to the simulation calculation.

  The initial input model of the nanoparticles and the input model of the intermediate nanoparticles are composed of various types of lattice planes. Since it is composed of various types of lattice planes, it is difficult to illustrate an input model that is close to reality.

  Although it is difficult to say that there are many types, for example, as an input model, it is assumed that a finite periodic part is formed by a simple cubic lattice, and three planes of {100}, {110}, and {111} are exposed. Assuming that the particles are substantially equal in length in each direction of the nanoparticles, a structure model of the nanoparticles as shown in FIG. 9B is obtained.

  If the partial atomic coordinate method of each atom as shown in the second embodiment is applied to this nanoparticle model, the existing position of each atom can be reproduced as shown in FIG. 9 (c).

  As described above, when n electron diffraction patterns are obtained in an experiment, by applying the present embodiment, it is possible to estimate not only the shape of nanoparticles composed of various kinds of surfaces but also the position of each atom. It becomes possible.

(Seventh embodiment)
Next, a seventh embodiment of the microscopic analysis method for fine particles according to the present invention will be described.
The seventh embodiment focuses on the intensity of the main spot and Laue sub spot.

From the case where the structural change of the finite periodic portion of the nanoparticles to be measured is small to the case where the structural change is large, it is described sequentially.
FIG. 10 is a schematic diagram (a) and graph diagrams (be) according to the eighth embodiment of the present invention.

First, a case where the change in the finite periodic portion of the nanoparticle to be measured is extremely small will be described.
For example, consider a case in which the crystal system remains a cubic lattice and the point group classification in the space group is m-3m, but the space group itself has changed. In order to confirm such a small change, attention should be paid to the intensity of the main spot.

FIG. 10 shows a case where the space group changes from Fm-3m to Fd-3m. This change corresponds to a case where atoms are alternately shifted along a diagonal line in a cubic system.
For example, when an electron diffraction pattern is photographed at [101] incidence, as shown in FIG. 10A, it is difficult to find a change from Fm-3m to Fd-3m due to the multiple scattering inherent in the transmission electron microscope. .

Therefore, when the sample is tilted around the [010] axis, as shown in FIG. 10B, both the [020] and [0-20] spots of Fm-3m remain on the electron diffraction pattern. , Fd-3m, [020] and [0-20] spots disappear.
As described above, by confirming the disappearance of the spot, it is possible to determine the shape of the nanoparticles after capturing the atomic displacement that does not change the point group classification although the space group has changed.

Next, an embodiment focusing on the strength change of the Laue subpot will be described.
First, the intensities of the main spot and Laue sub-spot of the obtained electron diffraction pattern are obtained. The intensities of the main spot and Laue subspot can also be calculated using a real grid-reciprocal grid transform.
However, in order to estimate the intensity with high accuracy from observation data, it is desirable to use the Fourier transform-Inverse Fourier transform as described in the second embodiment.

  Since each atom of the finite periodic portion of the nanoparticle is a nanoparticle, the partial coordinates may be displaced from the ideal position. FIG. 11A is a graph schematically showing a change in the intensity of the Laue sub spot due to the displacement.

  When the displacement of each atom occurs, the optical path difference of the electron beam varies due to the different atoms. Therefore, in general, the interference effect is reduced to be diffuse, and the intensity of each sub-spot is reduced. The intensity of each Laue subspot is affected by diffraction generated from all grating planes.

Therefore, the spot intensity increase / decrease amount of each Laue subspot shown in (1) to (3) in FIG. 11A differs for each peak according to the displacement amount of each atom.
Therefore, if the intensity analysis is performed by using the Fourier transform, it is possible to estimate the position (displacement amount) of each atom in each lattice plane.

  Usually, the Laue spot intensity is (1) = (1) ′, (2) = (2) ′, (3) = (3) ′, etc., but if the center of symmetry itself is shifted, ( 1) ≒ (1) ′, (2) ≒ (2) ′, (3) ≒ (3) ′.

  In the conventional method of focusing on the main spot and ignoring the Laue sub spot, the spot intensity of the main spot is extremely strong, so it is extremely difficult to detect a slight decrease in the main spot intensity. Therefore, it is extremely unlikely that the displacement of the atomic position and the change of the space group due to the displacement are captured.

  On the other hand, if attention is paid to the intensity change of the Laue sub-spot as in the present embodiment, it is possible to analyze the amount of position fluctuation of atoms in the periodic part by utilizing the fact that nanoparticles have finite periodicity. .

Next, a case where the present embodiment is applied to a solid solution will be described.
Further, the analysis of the solid solution state of a solid solution is extremely important, and there are various solid solutions ranging from metal materials that have been around for a long time to battery materials that have recently attracted attention.

  The analysis of the solid volume state may be made, for example, in the case of a full-rate solid alloy composed of two types of elements A and B, assuming that the alloy is composed of a single atom if it is not ordered.

  If the two types of elements A and B are ordered, atoms will be alternately arranged as ABAB. Then, the crystal lattice has a double period, and the magnitude of the scattering vector is halved. Therefore, as shown in FIG. 11B, the main spot having a half period with respect to the period of FIG. (Satellite spot).

  However, it may be difficult to observe the satellite spot because the atomic numbers of the two types of atoms are close to each other. Even in such a case, if the microanalyzer according to the present invention is applied, it is possible to know the degree of ordering and the plane orientation from the Laue subspot array and the intensity by ordering, as shown in the box of FIG. Become.

  FIG. 11C shows the intensity of the Laue sub-spot when the alloy has a concentration gradient. When there is a concentration gradient in the alloy state, the intensity of the Laue sub-spot changes, and the peak intensity of each Laue sub-spot increases and decreases as indicated by the arrows shown in FIG. Since the method of increasing and decreasing the peak intensity depends on the combination of atoms and cannot be unambiguously indicated, the increase and decrease indicated by arrows in FIG. 11C is for convenience of explanation of the embodiment.

  Since the Laue sub-spot is very weak and the peak intensity increases and decreases very little, it is effective to apply the Fourier transform-Inverse Fourier transform to accurately estimate the spot intensity. By analyzing the peak intensity of the Laue sub spot, it becomes possible to know the alloy concentration gradient of the finite periodic portion for each plane orientation.

The case where the present embodiment is applied to a clustered and amorphous sample will be described.
In the clustered nanoparticles, the distance between the atom of interest and the atom closest to the atom is almost the same as that of the crystallized nanoparticle, but the orientation is shifted. Further, the arrangement of atoms farther than the atom closest to the target atom is in an irregular state when viewed from the target atom. Also in this case, since it depends on the combination of atoms and the coordination method, the application of Fourier transform-Inverse Fourier transform is effective for analyzing the increase / decrease of the peak intensity of the Laue subspot.

The change of the Laue sub-spot in the clustered nanoparticles appears in the increase and decrease of the intensity of the Laue sub-spot as shown by the solid line arrow in FIG.
Since the Laue subspot is generated with reference to the central atom of the cluster, the peak intensity of the Laue subspot increases and decreases while maintaining left-right symmetry with respect to the scattering vector (the origin of the horizontal axis of the graph) 0.
In addition, due to the disorder of the coordinates of each atom, many weaker spots than the Laue sub-spot appear between the peaks of the Laue sub-spot, specifically, at a portion surrounded by a solid line in FIG.

  The weak spot is expected to be observed as a disorder of the foot of the Laue sub spot, rather than a spot, depending on observation accuracy. However, it is possible to estimate the degree of clustering of nanoparticles for each plane orientation by reproducing this disturbance by simulation using Fourier transform-Inverse Fourier transform.

  When the degree of disorder of the nanoparticles further advances and reaches a state in which the nanoparticles can no longer be considered as an amorphous state beyond the cluster state, the main spot intensity decreases as indicated by a broken arrow in FIG. 11D.

  The main spot (not shown, but the main spot which is located at the position of the horizontal axis at 6, 9,...) Away from the scattering vector (origin of the horizontal axis of the graph) 0, the further away from the origin, The spot intensity decreases.

  The completely amorphized nanoparticles are the first main spot next to the main spot having a value of 0 on the horizontal axis corresponding to the atom of interest (the main spot having a value of 3 and -3 on the horizontal axis). ) Is substantially zero.

Similar behavior is also reflected in the half width of the main spot, as indicated by the dashed arrow in FIG. 11D, and the half width becomes larger as the amorphization proceeds.
Therefore, it is possible to estimate the degree of amorphization for each lattice plane orientation by analyzing the spot intensity and the half-value width using Fourier transform-Inverse Fourier transform.

(Eighth embodiment)
Next, an eighth embodiment of the microscopic analysis method for fine particles according to the present invention will be described.
In the eighth embodiment, a plurality of electron diffraction patterns are obtained by changing the acceleration voltage of a parallel beam irradiated in a beam irradiation step. Spots and Laue subspots can be obtained.

  FIG. 12 is a graph according to the eighth embodiment of the present invention. By sequentially decreasing the acceleration voltage, the radius of the Ewald sphere created by the electron beam transmitted through the nanoparticles 161 to be measured is reduced.

  When the acceleration voltage is high and the diameter of the Ewald sphere is large like the Ewald sphere 162, a large number of main spots and their Laue sub-spots can be obtained as shown by black circles in FIG.

  Next, assuming that the diameter of the Ewald sphere is slightly smaller than that of the Ewald sphere 162 like the Ewald sphere 163 by lowering the acceleration voltage, as shown by a hatched circle in FIG. The main spot and its Laue subspot can be obtained.

Similarly, when the acceleration voltage is further reduced to make the diameter of the Ewald sphere smaller than that of the Ewald sphere 163, as shown by a check circle in FIG. 12, another type of main spot and its Laue subspot are obtained. be able to.
As described above, by further changing the diameter of the Ewald sphere, the main spot and its Laue sub-spot can be obtained in various directions.

  Since the present embodiment does not involve the sample tilt, the number of steps for setting the sample and the like is reduced to substantially the same level as the number of steps for taking one electron diffraction pattern, and then the number of electron diffraction patterns is reduced. As in the case where the image is taken, a main spot and its Laue sub-spots for a number of directions are obtained.

  Further, since the sample is not tilted, the arrangement and intensity of the main spot and its Laue sub spot can be obtained without being involved in mechanical error factors such as an eccentric error when the sample is tilted.

(Automation)
In a situation where the above-described analysis is performed on a large number of electronic analysis figures exceeding 100, it is difficult to perform the analysis using only human power. Furthermore, when performing the Fourier transform-inverse Fourier transform as shown in the second embodiment, it takes an enormous amount of time for calculation.
In such a situation, an electron microscope photographing system in which photographing of an electron diffraction pattern is automated is effective.

By applying the automated electron microscopy imaging system to the step of clarifying the structure of the finite periodic part and the repeated calculation of Fourier transform-Fourier inverse transform simulation, where various high-order diffraction spots appear Can be predicted.
Then, the Tilt and Azimuth angles of the electron microscope are moved to the positions where the predicted diffraction spots appear.

  Since the Tilt and Azimuth angles are usually slightly different from the designated position due to eccentricity error or the like, the Tilt and Azimuth angles are changed at the stage where this movement is completed by using the increase or decrease of the spot intensity of the focused index. Fine tune.

  At the stage where the above fine adjustment is completed, a target electron diffraction pattern is photographed. Embodiments 1 to 8 are applied to the obtained electron diffraction. By doing so, photographing of more than 100 electron diffraction patterns becomes realistic.

  It is also important to determine as soon as possible whether the obtained electron diffraction pattern is appropriate. To this end, it is desirable to connect a computer to the automated electron microscope imaging system and immediately perform Fourier-Fourier inverse transformation simulation.

  As a result, when it is found that the azimuth correction at the time of capturing the electron diffraction pattern is insufficient, the electron diffraction pattern is re-photographed. In order to smoothly perform the Fourier-Fourier inverse transform simulation, it is preferable to use a supercomputer (HPC: High Performance Computer).

Hereinafter, the contents of the present invention will be described more specifically with reference to examples.
An electron diffraction pattern was photographed with respect to platinum (Pt) particles with the incident direction of the electron beam set to [100]. FIG. 13 is a photograph showing an example, and FIG. 14 is a photograph showing a comparative example.
FIG. 13 shows an image obtained by narrowing the opening angle of the electron beam to about 1 mrad and the spot diameter to about 5 nm by using a converging lens, and irradiating this with nanoparticles.
On the other hand, FIG. 14 shows an example in which an electron beam having an opening angle of about 0.1 mrad and a spot diameter of about 50 nm is applied to the same nanoparticles as in FIG. It was taken.

  As can be seen from the inserted photograph at the lower right of FIG. 13, the size of the target Pt nanoparticle is 2 to 4 nm. Pt nanoparticles are supported on carbon. In practicing the present invention, the spot diameter of the electron beam irradiated to the nanoparticles is important. However, it is difficult to form a nanobeam electron beam having a small diameter as described in each of the above embodiments with the current electron microscope.

Therefore, the photograph of the example shown in FIG. 13 is not ideal because it is irradiated with an electron beam having a spot diameter larger than that of the nanoparticles. Nevertheless, it can be seen that an electron diffraction pattern with clearly clear diffraction spots is obtained as compared with the photograph of the comparative example shown in FIG.
Needless to say, if an ideal small-diameter nano-beam electron beam having a spot diameter substantially equal to the diameter of the nanoparticles as shown in each of the above-described embodiments is applied, the supported carbon shown in FIG. Spots (obstruction spots) due to this can be greatly suppressed.

  FIG. 15 shows a spectrum diagram in the direction of the arrow shown in the photograph of the above example. As shown in the spectrum diagram, the (111) lattice spacing obtained from the value of the scattering vector was 2.2 °. It almost coincided with the literature value of the (111) lattice spacing of Pt.

  Further, as shown in FIG. 16, the reciprocal value of the scattering vector of the Laue subpeak was 29 °. Therefore, it can be seen that the Pt nanoparticles consist of 29 / 2.2２13 lattice planes. That is, this example shows that the length of the Pt nanoparticles in the [111] direction is 29 °. Similar analysis is performed for various directions. Thereby, the length of the nanoparticles in substantially all directions can be known.

  Pt belongs to the Fm-3m (FCC) structure. Since the (111) lattice plane is constituted by ABC stacking, (13-1) × 3 = 36 Pt atoms are arranged in the [111] direction.

  The microscopic analysis method for fine particles of the present invention and the analyzer used for the microscopic analysis method for fine particles can contribute to the development and analysis of materials exhibiting characteristics by being nanoparticles, such as photocatalysts and fuel cell catalysts.

101 Main Spot 102 Laue Sub Spot 111 Sample 112 Convergent Lens 113 First Intermediate Lens 114 Second Intermediate Lens 115 Projection Lens 116 Projection Surface 117 0th-Order Reflection Mechanism 118 Main Spot Cutoff Mechanism 119 Detector 120 Restricted Field Stop 161 Nanoparticle 162 Ebalt Sphere (large)
163 Ewald Ball (Medium)
164 Ewald Ball (Small)

Claims (8)

A beam irradiation step of irradiating the parallel beam to the fine particles,
A detection step of photographing a plurality of electron diffraction patterns by changing the photographing angle and / or the size of the Ewald sphere for the fine particles;
An array confirmation step of analyzing the diffraction spot of the electron diffraction pattern to determine the shape of the fine particles,
Including
The arrangement confirming step compares an input model predicting the size and shape of the fine particles with an array of Laue sub spots of the plurality of electron diffraction patterns, and determines the shape of the fine particles by correcting the input model. A method for microscopic analysis of fine particles, characterized in that: The microparticle analysis method according to claim 1, wherein the fine particles have a cluster portion and / or an amorphous portion. 3. The method for microscopic analysis of fine particles according to claim 1, wherein the arrangement confirming step compares an input model with an arrangement of main spots and Laue sub-spots of the plurality of electron diffraction patterns.   The method according to any one of claims 1 to 3, wherein the input model is constructed using a lattice constant of a finite periodic part, a space group, and partial atomic coordinates. The sequence confirmation step includes an atomic position determination process,
  The method for microscopic analysis of fine particles according to any one of claims 1 to 4, wherein the atomic position determination processing determines an atomic position using Fourier transform and inverse Fourier transform. The microscopic analysis method for fine particles according to any one of claims 1 to 5, wherein the sequence confirmation step includes an analysis of Laue subspot intensity. 7. The microscopic microscope according to claim 1, wherein an electron diffraction pattern is photographed with a parallel electron beam system after confirming a photographing angle of the fine particle with a convergent electron diffraction system. Analysis method.   An electron beam irradiator that irradiates the sample with a parallel electron beam, an intermediate lens, a projection lens, a zero-order reflection blocking mechanism that blocks zero-order reflection, a main spot blocking mechanism, and detection that detects a Laue sub-spot of the sample. With a container,
  A microscopic analyzer for fine particles, characterized in that the electron beam irradiator has a converging lens for converging an electron beam, and the converging lens can form an opening angle of less than 1.0 mrad.

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