There are two ways to measure stresses by means of X-rays, namely the counter method and the photographical method. There is tendency in recent years that the counter method has been gradually being employed in laboratories and factories because of its affording possibilities of easy handling and quick measuring. It is well known, however, that the photographical method has hither to been more frequently applied than the counter method to the measurement of residual stress on local and small positions. There are several difficult problems about the measuring technique of the conventional photographical method, for example, the difficulty in powdering the specimen surface with calibrating material and the difficulty in determination of peak positition on the broadened diffraction lines, since the distance between the diffraction lines from the specimen the calibrating material is close on the film.
In the present study a new method of determining X-ray stress from this point of view by the photographical method without using calibrating material has been tested and those features are mentioned. The distance between the X-ray film and the surface of specimen can be precisely determined by the diameter of diffraction ring of the calibrating material which powders the specimen surface by the conventional method. Bragg's angle of specimen is obtained by using the value of the distance and the radius of the diffraction ring of the specimen. On the other hand, by a new method which is called the double exposure method, the first exposure is given on the X-ray film from which the distance to the specimen is l₁, and when the distance between the specimen is set at l₂, the second diffraction ring of the specimen is photographed. Bragg's angle 2θ is calculated from Δl, the difference between the distance l₁ and the distance l₂, and Δr, the difference between the radius r of the diffraction rings at the first and second exposures, as 2θ=tan^-1[-Δr/Δl₁].
As was expected, it has been made clear from the theoretical and the experimental discussions that this double exposure method has many advantages as follows:
(1) If the difference between the radius of both the diffraction rings, Δr, is larger than a quarter of diameter of the diffraction ring of the specimen on the conventional method, the accuracy of the measured value of stress by the double exposure method is better than that by the conventional method.
(2) The two diffractions forming the similar patterns, the distance between the peaks of diffraction intensity curves, have been easily measured.
(3) Irradiation of X-ray beams can be easily pointed at the measuring position on the specimen without using the metallic powder as calibrating materials.
(4) There is a difficult problem in measuring by the conventional method, the stress on the specimen with broad diffraction lines, because these diffraction lines of the specimen interface with those of the calibrating material. In contrast with above mentioned case, the stress measurement of such a specimen is made possible by employing this double exposure method.

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A new type X-ray diffraction apparatus has been developed to make measurement of stress in machine parts and their structure with ease. It is especially devised to make it practicable to make measurement of samples which show broad or spotty diffraction rings.
The distance of X-rays from the source to the sample surface to the detector of the stress measurement goniometer is shortened as much as possible. A full wave rectification system with smoothing condenser is used as X-ray generator, so that the intensity of diffracted X-rays may be increased from four to five times as much as the old type apparatus. The scanning range of the detector is broadened to 140°∼170°in diffraction angle 2θ. The counting range, time constant, scanning speed of detector and speed of chart feed can be chosen widely according to the intensity and broadening of diffraction lines.
Consequently, the data of broad diffraction lines can be processed easily, and the accuracy of the measurement is improved. The scanning speed of the detector for the measurement of diffraction lines with high X-ray intensity becomes four times as fast as the old type apparatus.
The local stress of the sample with coarse grains can be measured without oscillating the sample, because incidental X-rays can be oscillated around the irradiated point. Without removing the sample, the stress can be measured also by θ-2θ method.
The apparatus is applicable to automatic measurement or automatic digital data processing.

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There have been a number of methods proposed, of calculating the elastic constants of polycrystalline material from those of its constituent crystals. As the most fundamental boundary conditions of these methods, Reuß proposed an assumption that stress is uniform throughout the aggregate. On the other hand, Voigt proposed another assumption that it is strain that is uniform. They treated only those cases in which the constituent crystals are randomly oriented. However, most polycrystalline metals contain crystals with lattice orientations that are not random, but instead are aligned to some degree about a particular orientation or about a set of orientations because of the mechanical and thermal history of the metals. The elastic constants of a single crystal are generally anisotropic, and therefore, the elastic constants of a polycrystalline metal should be influenced by its preferred orientation or texture.
In the present paper, a method of estimating the bulk elastic constants of a polycrystalline metal with fiber texture is proposed, and the calculated results of a polycrystalline α-iron with [110] texture are presented as an example. Two kinds of bulk elastic constants are treated here. One is termed the mechanical elastic constant which is the constant averaged over all the crystals constituting polycrystalline aggregate. The other is termed the X-ray elastic constant which is the constant averaged over only the select crystals that are in the correct direction to satisfy Bragg's condition.
It is found from the results of α-iron that the X-ray elastic constants calculated on the basis of the Reuß condition are reduced about 20 percent by the development of [110] fiber texture. On the other hand, neither the X-ray elastic constants based on the Voigt condition nor the mechanical elastic constants are estimated to be decreased more than 5 percent by fibering.

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It is well known that the X-ray stress measurement is a unique and effective method of measuring stress, both micro and macro, in polycrystalline metals. So it is applied in wide fields of material engineering studies. There are a few problems, however, regarding the fundamentals of X-ray stress measurement. The metallic materials in practical use are polycrystalline, and in the X-ray stress measurement it is the average intraplane lattice strain, with a certain number of crystals favorably oriented with respect to the radiated characteristic X-rays, that is measured. Consequently, it is considered that the lattice strains observed by X-ray diffraction are closely related to the elastic or plastic anisotropy of the crystal grains, the direction of stress and the complex deformation mechanism.
On the effect of crystal anisotropy on the X-ray stress measurement reports of studies have so far been made by several investigators, and it is pointed out that the stress value differs in its measurement from one lattice plane to another on which the X-ray diffraction line is reflected. Though discussions on these problems have thus been made to a certain extent of X-ray measurement of the lattice strains, and of its dependence on the diffraction plane of the crystal, its details remain vague yet.
In order to solse these problems, the authors have attempted certain theoretical treatment based on six equations of generalized Hooke's law on the crystal elasticity, and compared them with experimental results of crystal plane dependence of lattice strains which was obtained from the X-ray stress measurement of low carbon steel, using (310), (211) and (220) diffraction lines. Detailed accounts of the experiments and the conclusion of this study are to be comprehensively given in the following paper.

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In the previous paper, the changes in the lattice strain and the line broadening on the diffraction line were analytically discussed, in view of the elastic anisotropy of crystal, in connection with the mechanism of elastic deformation. It was the aim of the discussion to make clear the relation between the results of the theoretical analysis and the experiments.
Three kinds of materials were used in these experiments, that is, annealed plate specimens of 0.12 percent carbon steel, commercial pure aluminium and copper. All the specimens were electro polished before being exposed to X-rays. The characteristic X-rays of Co Kα, Cr Kα, Fe Kα, and Cu Kα were radiated on the specimen surface through parallel beam slit of divergent angle of 0.25 degree, and the strains were measured by using the diffraction from (422), (420), (400), (331), (310), (220) and (211) atomic planes. The specimens were stressed stepwise by the tensile testing machine, and at several stages of applied stress, the X-ray beams were radiated to the center of the specimen surface in vertical and oblique incidence with several anglesψ. The strain was measured by the conventional sin²ψ method using the counter technique. The value of the lattice strain (ε_ψ) was calculated from the measurement of the diffraction angle of intensity distribution curve by using automatic recorder. From the slope of the lattice strain ε_ψ-σ_m diagram for several applied stresses the ε/σ-sin²ψ curve was drawn by using the method of least square. From these slopes the elastic constant was calculated for each diffraction plane.
The conclusions of the present study are as follows.
(1) The presence of diffraction plane dependence is considered to have been caused by elastic anisotropy. In view of the fact, however, that certain metals contain texture, it is considered to be necessary that the effect of the texture will be taken into consideration in addition to the above mentioned analytical treatment.
(2) In connection with this study, the authors measured the integral breadth in the stage of elastic defomation by the X-ray method, and compared it with the result of analysis of dispersion of micro strain in the stage. The diffraction plane dependence is presented in this case also.

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In making X-ray measurements of materials under known uniaxial elastic stress, we can determine the X-ray elastic constants S₁ and S₂/2 experimentally. Of the same material, however, the measured values of the X-ray elastic constants usually vary with the wavelength of the radiation used, or the Miller indices of the measured lattice planes. Here, we call this the diffraction plane dependency of X-ray elastic constants.
The diffraction plane dependency of X-ray elastic constants has hitherto been discussed by many investigators analytically and experimentally. Many of them have used the different characteristic radiation in order to catch the diffraction line on the different lattice plane. But it is more reasonable to use the same characteristic radiation in the study of the diffraction plane dependency.
In this paper, the strain determination of different lattice planes was made by means of specially constructed devise attached to the X-ray diffractometer which made it possible to use the sin²ψ method not only at a high diffraction angle, but also at a very low diffraction angle, and so made it possible to use the same characteristic radiation for the study of the diffraction plane dependency of X-ray elastic constants.
In these experiments, which were carried out under uniaxial tension at various elastic stress levels, the 0.2% low-carbon steel specimens were used. The mechanism of the stress measurements is indicated in Fig. 2. The characteristic radiations and diffraction planes used are as follows; Cr Kα radiation-(110) (200) (211), FeKα radiation-(110) (200) (211) (220), CoKα radiation-(200) (211) (220) (310).
The experiments show that the measured X-ray elastic constants S₁ and S₂/2 have no relation to the wavelength of the radiation used, but depend upon the measured (hkl) lattice plane notably; that is, the absolute values of S₁ and S₂/2 on (110) and (211) lattice planes are the same and are lower than that on a (310) lattice plane. The absolute value of S₁ and S₂/2 on a (200) lattice plane is the largest, and the X-ray elastic constants have linear relation to the orientation function.

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X-ray diffraction patterns of polycrystal specimens are generally spotted. When we are to determine the residual stress of such specimens by X-rays, the specimens must be moved to produce a continuous series of diffraction rings. Using annealed low carbon steel plates, we experimentally investigated how their diffraction rings appeared under horizontal and rotative movements.
For the experiment we have used Cr Kα line of parallel beam X-rays and divergence X-rays. The X-ray diffraction patterns obtained from this experiment were invariably spotted ones extended in a specific direction which depends on the manner of the movement of the specimen. This direction agreed with the specimen movement in the case of horizontal movement, and with the direction normal to the rotative axis in the case of rotative movement.
According to the observation, the mechanism of variation in the spotted pattern differ from case to case.
Though the geometric diffractive conditions of the X-rays and specimens are the same, it has been revealed that the difference in mechanism due to the state of crystal cohesion and that due to the solar slit scattering of parallel beam X-rays are superposed upon each other.
Moreover, assuming that the X-ray diffraction pattern is caused by the Bragg reflection, we have theoretically established the conditions for diffraction ring to separate itself into Kα₁ line and Kα₂ line, with respect to the parallel beam X-rays and the divergence X-rays. The condition thus obtained agreed well with the experimental results.
In the X-ray diffraction patterns obtained from the experiment, a diffraction pattern for the satellite line of Kα line could be recognized.

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The diffraction intensity of Debye-Scherrer ring is generally homogeneous irrespective of the direction in which the incident beams shoot, when the crystallites are distributed in all directions in random orientation. There are some powder samples, however, for instance, those having tendency to cleave, of which the intensity distribution depends largely on the way how to mount them on the specimen holder. The present study was designed to prepare a special specimen of this kind of substance for the counter diffractometer.
In the process of packing powder specimen on the specimen holder preference can take place in orientation in the formation of crystallites, and this can be prevented by jagged specimen surface. In view of this fact special mould for making the specimen was devised. Fig. 4 shows the schematic representation of this. The final specimen was produced by stamping the powder by the use of this mould.
The special mould was applied to the analysis of the powder pattern of PbO and β-alumina whose intensities are generally uneven due to preference taking place in orientation during the packing in the formation of crystallites.
The details of the experiment are listed in Table I∼IV, in which it is clearly seen that the homogeneous powder patterns can be obtained by the choice of suitable cycle of jaggedness to be produced on the surface of the specimen.

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The authors previously reported on the behaviour of residual stress and change in microstructure due to cyclic stressing by rotary bending of tufftrided carbon steel specimens. They also pointed out that no change was observed in residual stress and microstructure on the surface layer of the specimen during the fatigue process. The strength of tufftrided steel was influenced by the amount of precipitation of nitriding materials in the matrix.
In order to clarify the improved mechanism of fatigue-resistance of tufftrided steel, however, it is required that more detailed investigation will be made. It is especially important to make inquiry of the effects of mean stress superposed to the cyclic stresses and the contribution of residual stress. Investigation was therefore made of the fatigue limit and the fatigue life on the alternate bending with mean stress, and of the change in residual stress during the fatigue process.
Plate type 0.07% carbon steel specimens were used as the experimental material, and these specimens, after being machined and annealed, were tufftrided at 570°C for 90min. The alternate stress and mean stress were applied to the specimen by Schenck type reversed bending fatigue testing machine. The residual stress in the specimen surface on the way of fatigue process were measured by X-rays. By taking the residual stress as a mean stress and by calculating the equivalent stress amplitude in reversed bending, the contributions of residual stress to the fatigue limit and the fatigue life were studied. An ideal S-N curve, which did not include the influence of residual stress, was introduced.
The results of the experiment and discussions thereanent are summarized as follows:
(1) Both the fatigue limit and the fatigue life decrease as the mean stress increases. The fatigue limit in the case of supeprosed mean stress coincides with the initiation of plastic deformation.
(2) The compressive residual stress in the specimen surface, where the tensile mean stress is loaded with cyclic stress, almost disappears in the initial stage of fatigue life. On the other hand, the residual stress on the compression side slightly decreases at about-20kg/mm² in the initial stage, but this value remains to the fracture. It is considered that there is no contribution of compressive residual stress of the fatigue strength.
(3) Based on the ideal S-N curve, the contribution of residual stress to fatigue limit which is improved by tufftriding is about 14∼30% in case the fatigue is caused by reverse bending.
On the other hand, it is evident that 50∼90% of the whole fatigue life is influenced by the compressive residual stress.

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The relation between the strength of surface layer measured by X-rays under static bending and the fatigue strength was investigated on various iron castings.
Fine flake graphite cast iron, malleable cast iron and spheroidal graphite cast iron with ferritic, sorbitic and pearlitic matrix were used for the materials.
The nominal bending stress and the X-ray stress of iron castings were related linearly in the early steps of loading. Then passing through some non-linear range, the values of X-ray stress became constant. The nominal bending stress, where the relationship ceased to be linear, were designated the proportional limit of the surface layer σ_eX, and the nominal stress, where the X-ray stress became constant, was designated the yield stress of the surface layer σ_sX.
The endurance limit σ_w' of iron castings was confirmed by the rotating bending fatigue test.
The yield stress σ_sM and rotating bending endurance limit σ'_wM of steels, which corresponded to the matrix of iron castings, was approximately estimated from the micro-structure and hardness of the matrix.
σ'_wM/σ'_w which represented the ratio of endurance limit of steel and iron casting was defined as fatigue notch factor by graphite β_g. σ_sM/σ_eX or σ_sM/σ_sX corresponded to the ratio of the yield strength of matrix to the statical strength of surface layer of iron casting.
The fatigue notch factor by graphite β_g showed a linear relation to the ratio σ_sM/σ_eX or σ_sM/σ_sX and fairly agreed to the latter in value, despite the difference in the micro-structure or hardness of the iron castings tested.
Consequently, the endurance limit of iron castings could be obtained from estimation of the following expression.
σ_w'=σ'_wM/β_g=σ'_wM/σ_sM·σ_sX.

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From several reports of scientific experiments that have been recorded in the literature of science it appears that the practice of removing layer after layer of the specimen surface during the low-cycle fatigue testing has an effect of lengthening the fatigue life, and it is surmised as possible that this phenomenon is due to some microstructural difference that is created by cyclic straining between the surface and the inside of the specimen. And detecting of this structural difference would provide more resources to the study on low-cycle fatigue behavior. From this point of view, the authors applied the X-ray diffraction techniques to observation of the microstructural change in the surface and the inside of the specimen during constant strain-cycling.
The following facts have been revealed as the result of this investigation:
(1) It is revealed by the X-ray microbeam diffraction technique that there is subgrain formation and development not only on the surface layer but also inside the specimen in the process of low-cycle fatigue before visible cracks occur.
(2) The period of the low-cycle fatigue process from distribution of the integral breadth of its X-ray diffraction peaks from the surface in the direction of depth to the occurrence of visible cracks is divided into two stages. The first stage marks larger increase, with increase in the number of strain cycling, in the integral breadth on th surface of the specimen than inside. During the second stage the integral breadth tends to decrease on the surface of the specimen, but keeps increasing inside. The integral breadth shows smaller value on the surface than inside around the failure point.
(3) Microscopic cracks are found to occur toward the end of the first stage, and develop in the direction of depth from the surface during the second stage.
(4) The average value of integral breadth obtained from inside the specimen has a linear relation with the stress range during the low-cycle fatigue before fracture.

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The 0.45% carbon steel specimens, of which surface layers were perfectly quenched in various depth by induction hardening, were examined in the rotating bending fatigue test machine to ascertain the relation between fatigue endurance limit and residual stress. The result shows that both hardness pattern and residual stress together are not sufficient to be adequate to clarify the interrelations among the endurance limit, hardened depth and residual stress. Some more factors are required for the purpose. Therefore investigations must be made in that line.

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Recently transmission electron microscopy has come to be widely used as efficient means to study the dislocation microstructural changes during the fatigue of f.c.c. metals, but not so widely in the case of b.c.c. metals, because of difficulties in dealing with the thin foil technique in the latter instance. No examination has so far been made of the fatigue behaviors of cold worked b.c.c. metals, such as rolled low carbon steel, by means of transmission electron microscope, inspite of the important practical problems involved therein.
In this paper report is made of the direct observation of dislocation microstructure conducted by means of transmission electron microscope during the fatigue, with a view to clarifying the fatigue deformation mechanism of the annealed and the rolled low carbon steels, with discussion on the strengthening mechanism in the fatigue strength of the rolled material.
The results are summarized as follows:
(1) In the case of the annealed material, the initial stage in fatigue deformation process is characterized by increase in dislocation density, arrays of dislocation loops in parallel rows, dislocations in clusters and tilt boundaries composed of dislocation network. The second stage marks formation of substructure, and the final stage is reached by fragmentation of subgrains. This subgrain size is 2∼4μ in diameter.
(2) The rolled material contains, before it is subjected to fatigue test, various lattice defects which cause imperfect substructures in its crystal grains. These imperfect substructures can be divided into two main regions; one is the rough substructure region in which the difference in dislocation density between the sub-boundary and the inside of the subgrain is low, and the dislocation density comparatively low on the whole, and the other is the band structure region which is composed of lamellar distribution of highly tangled dislocation band. This band structure region presents almost no change during the fatigue. In that rough substructure region, there have been, with increase in the number of cycles, interaction of dislocations, annihilation of dislocation, and reversible movement of dislocations toward the positions of lower potential energy, and then there have been formed distinct substructures to a larger number than at the time of rolling process. This subgrain size is 2∼6μ in diameter.
(3) The fatigue life of the rolled material is prolonged longer than in the case of the annealed one. This phenomenon is due to the resistance to the mobility of dislocation under the cyclic loading of the small dislocation loops, dislocation network and these tangled dislocations in the rolled material. This fact shows that greater energy is required for the formation of favorable substructures for crack initiation in the rolled material than those in the annealed one.

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Three kinds of low carbon steel were examined to study the relation between the applied stress condition and the crack propagating rate. The respectively chosen specimens were those initially annealed, those stretched and those recovered. These were all subjected to alternating push-pull load.
The study was made from the view point of micro-structural change at the crack tip observed by the X-ray microbeam diffraction technique.

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In order to make clear the mechanism of the fatigue crack propagation, it is important to find out what sort of stress it is at the tip of the crack that will start again when a cracked specimen is subjected to repeated loading. It may be considered that there is unstable equilibrium in the strength of the material at the tip of the crack, and that the applied stress determines whether the crack will propagate itself or not.
There are many factors that affect the strength of the material ahead the tip of the crack. These factors, however, may be divided into two groups, the structural factors and the mechanical factors. The former group consists of the crystal forms, grain size, behaviours of the dislocation, point defects and behaviours of solute atom (Impurities), etc. The latter is defined by the interrelation between the external stress and the distribution of the internal stress (Residual stress).
Therefore, to clarify the mechanism of the crack propagation from the point of view of the mechanical factor, the residual stresses at the tip of the crack must be measured, first of all. But, hitherto, as the stress measurement technique of a localized region such as the crack tip has not been established, it seems that no investigation has so far been made yet on this line.
After considerations on these points, in the present study, the authors measured the microscopic and macroscopic residual stresses in the crystals which exist at the tip of 12 sorts of the fatigue crack in annealed copper specimens, respectively, with aids of flow stress curve during the uniaxial stretching process and the crystal oscillation X-ray microbeam diffraction photographic technique.
The results are as follows.
(1) The microscopic and macroscopic residual stresses measured in the crystals ahead the tip of the crack increased with increase in the crack length, while the minimum stress required to propagate the fatigue crack σ_min.c decreased with the crack length.
(2) The ratios of stress concentration α_f.c and α_r.c which were confined by dividing the strength of the material at the tip of the crack by σ_min.c increased with the crack length.
However, it seems to be necessary to make further inquiry into the case of α_f.c since no arrangement of the dislocations were taken into consideration in the estimation of the strength of the material at the crack tip, though the density alone was duly considered.
(3) If the macroscopic residual stress is regarded as the mean stress, it may be considered that the stress state of the cyclic loading at the crack tip was compressed during the fatigue process. In this situation, however, it is clear experimentally that at the initiation of plastical deformation in the zone new crack appeared at the crack tip.

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In discussing the mechanism of fatigue crack propagation in metals, there are two important problems as follows: (1) how the plastic deformation and the stress distribution at crack tips are related to the applied gross stress and crack length, and (2) how propagation rate is controlled by the plastic deformation or by the stress distribution at crack tips.
In this study, as the first step on the way of investigation on the mechanism of fatigue crack growth, the plane specimens of coarse-grained pure iron with a shallow single edge notch were fatigued under completely reversed tension compression stress. The plastic deformation near the tip of the propagating cracks was examined by optical microscopy and the back-reflection X-ray microbeam Laue technique. The obtained results are summarized as follows:
(1) The slip bands zone size ahead of fatigue crack tips is about 15 percent of the plastic zone size predicted from Dugdale's model. The residual stress induced by fatigue will reduce the size of the plastic zone at the crack tips in cyclic stressing.
(2) The crack growth rate dl/dN is uniquely related to the slip bands zone ahead of the crack tips ξ, as
dl/dN=6.0×10^-5ξ^1.7,
regardless of the stress amplitude or the crack length. At the minimum rate for cracks to propagate, ξ takes the value of about 5μ. It is worthy of note that this result was obtained in the case of 15% cold-rolled 0.01% C iron.
(3) Where the slip bands are observed in the vicinity of cracks, the value of misorientation becomes larger, while the subgrain size becomes smaller. The misorientation β at the fatigue crack tip increases as crack growth rate dl/dN becomes larger and their relationship is expressed as,
dl/dN=2.2×10^-1β^3.5.

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The effects of cathodic charging of hydrogen in the solution of 1N-Na₂SO₄ were investigated by X-ray diffraction technique on the crystal structure, macro stress and micro strains in electrodeposited chromium obtained from chromic acid bath.
The results obtained are summarized as follows;
(1) It is found that electrodeposited b.c.c. chromium is partially transformed into h.c.p. type chromium hydride by hydrogen charged under the conditions of bath temperature range of 30 to 70°C, current densities of 1 to 10A/dm² and longer charging time.
(2) The lattice parameters on both (211) and (222) planes of electrodeposited b.c.c. chromium increase and the tensile stress on (211) plane measured by Sin²ψ method decrease with charging time. Further charging change the sign of macro stress but holding at room temperature reduce the effect of the charging on the macro stress. This behavior of macro stress under holding may be explained by the diffusion of hydrogen from matrix to surfaces, causing the expanded lattice to contract.
(3) The half-value breadths on both (211) and (222) planes of electrodeposited b.c.c. chromium increase with charging time. The increase in the diffraction line broadening was found as being caused by the fragmentation of crystals not so much by micro strains. However, large micro strains were measured in the transformed h.c.p. type chromium hydride.
(4) It is found that the effect of hydrogen charging on electrodeposited h.c.p. and f.c.c. type chromium hydride is not remarkable as it is detected by the line broadening as compared with the case of b.c.c. chromium.

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It is generally known that varieties of phenomena occur to austenitic stainless steel when hydrogen is introduced into it by the cathodic electrolytic method. Four sorts of austenitic stainless steels SUS 39, 27, 32 and 42 (Corresponds to AISI 301 304, 316 and 310) were examined. These specimens were prepared and hydrogen was introduced into them by the electrolytic method. The F.C.C. austenite was observed as tranceformed into B.C.C. Alpha and H.C.P. Epsilon, and at the same time, interstitial compounds of nickel were recognized. The effects of elastic stresses and prestraining were also measured

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On the impact defomation of metals, several studies have hitherto been reported.
In this study, the rotation of crystals and Laue asterisms caused by the application of static or impact loads to the coarse grained specimens of aluminium were observed by the Laue method.
The experiment was carried out on three sorts of crystals whose tension axes were nearly in the [110], [001] and [111] directions. On the other hand, shear strains in the slip plane [111] and rotations of slip direction [101] were calculated in both cases of static and impact loads.
The experimental result obtained may be summarized as follows. In each case of the crystal orientations, the rotation of crystals was larger in the static than in the impact loading. In both cases of the static and the impact loadings, the change of angle between the tension axis and the [112] direction was at the minimum with the [111] crystal. On the other hand, the result of calculation showed that the increase in shear strains was most remarkable in the [111] crystal in loadings. The Laue asterisms produced by the impact loading were more pronounced than those both generated by the static loading.

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