In gas turbines the rotor disks are highly stressed by centrifugal force caused by high speed rotation and are required to run for a long time at high temperature. So it is important for designers to predict the stress and the creep strain. Research on the creep of rotating disk, both theoretical and experimental, is indispensable for development of gas turbine and many theoretical papers have been published. But, as for experimental papers, the number is quite small, as the experiment necessitates complicated methods and is expensive and moreover requires time and efforts.
Wahl has developed an analytical study on steady state creep of rotating disks extending Bailey's method. In his study, creep rate at the given temperature is assumed to be a function of stress multiplied by a function of time i.e.ε=F(σ)f(t) and as a key expression between stress and creep rate the power function relation i.e. F(σ)=σⁿ is adopted. Two dimensional stress state is assumed because axial stress is thought negligible as compared with radial and tangential stresses which are constant across the thickness of the disk. Elastic and transient creep strains are assumed to be negligible as compared with the steady state creep strain. Calculations are carried out in the following three cases: (1) Mises criterion and Mises flow rule, (2) Tresca criterion and Mises flow rule and (3) Tresca criterion and its associated flow rule. The analytical results have been compared with the experimental results carried out by him and his collaborators. In the experiments, allegheny 418 stainless steel disks (12% Cr, 3%W) were tested at an uniform temperature of 1000°F (538°C), and the test time exceeded 900 hours. Disks were 12 inches (305mm) O.D., 2.5 inches (63.5mm) I.D. and 1 inch (25.4mm) in thickness. According to him, the analytical results obtained by means of Tresca criterion and its associated flow rule generally showed the closest agreement with the test data. Based upon this result he is continuing the analysis further.
Manson has presented an analytical method, in which effects due to elastic and plastic strains, disk contour and temperature gradient of disk are taken into account as well as the effect of creep strain. Deformation theory in plasticity is applied for the calculation of creep strain, and Mises criterion and flow rule are employed. Other assumptions are similar to those of Wahl. In numerical calculation the finite difference approach is used, but it is so complicated that the computation is almost impossible without electronic computors.
In this paper, the authors have adopted the assumptions similar to those of Wahl. Elastic and transient creep strains are assumed to be negligible as compared with steady state creep strains. Mises and Tresca criteria with Mises flow rule are applied in the calculation. In solving stress equations Runge-Kutta's method or finite difference approach is used. The results obtaind are close to those of Wahl, but the calculation procedure is simpler. Analytical results are compared with the test data obtained by means of hot spin tester manufactured by the authors. S 816 disk (41% Co, 20% Ni, 20% Cr) was tested and its sizes were 300mm O.D., 90mm I.D. and 20mm in thickness. Disk temperature was uniform over the disk and was 820°C. The results of the authors' analysis and experiments are compared with those of other researchers. Tresca criterion has given a good agreement with the authors' test data as it did in Wahl's experiment. However, further test data are necessary before a general conclusion can be drawn.

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We cast keel test blocks corresponding to 18-8 system austenite stainless steels used for large turbine castings-AISI 304, VIS 31, AISI 347, 316 Cb and 321-, and investigated mainly into the high temperature properties of these test blocks. The results obtained are as follows:
(1) Room temperature strength of the steel VIS 31 is slightly inferior to the others'. Ductility of the steel 304 is the highest, followed by the others' in the order of the steel VIS 31, 321, 347, and it is the lowest in the steel 316 Cb.
(2) Short time strength of the steel VIS 31 at 650°C is a little inferior to the others' and ductility of the steel 316 Cb at 650°C is the lowest among the tested steels.
(3) The ranking in rupture strength at 650°C is in the order of 347, 316 Cb, VIS 31, 321, and the steel 304 is the lowest.
As compared. mechanical properties of castings with properties of wrought steel materials corresponding to these cast steels, room temperature and short time high temperature strengths of castings are lower by about 5kg/mm² than those of wrought steel materials. Also ductilities of the former are about 10% inferior to those of the latter. Rupture strength at 650°C of the former is either equal to or a little inferior that of the latter. These castings pass through JIS corrosion resistant specification for wrought steel materials. When we compared the influence of water toughning or air cooling on the strength or the ductility in room and elevated temperatures of cast materials and corrosion resistant properties of theirs, we could not find any remarkable differences between them. Therefore, we could confirm the fact that the post-annealing of large and complicated castings as turbine castings permits of air cooling.

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Stress rupture tests with 40ton tester were conducted at 500 and 600°C for large welded specimens of stainless clad steel, Mn-Mo-Ni steel ASTM A 302 B clad with AISI 304L, which is essential to nuclear applications. The large specimen, 38mm thick and 40mm wide, showed much inferior property to usual small welded specimen.
The results of test at 600°C under a stress of 12kg/mm² are:
(1) The rupture time of transverse as-welded specimen was 55% of base metal.
(2) The rupture times of transverse welded specimens decreased through heat treatment after welding at 550, 650 and 870°C×1.5hr as low as to 40, 60 and 90%, respectively, of as-welded specimen.
(3) Machining the reinforcement flush with the surface of bese metal increased the rupture time of 550°C×1.5hr heat treated specimen by 1.5 times.
(4) Longitudinal welded specimen was approximately equivalent to base metal in rupture time.
(5) Transverse specimen manually welded showed about 50% longer rupture time than submerged-arc welded specimen.
Macroscopic and microscopic examinations showed that cracking in rupture occurred along bond in the graphitized heat-affected zone of base metal. The initiation of cracking occurred mostly at the intersection between the surface of bond of stainless weld metal and the cladding surface, viz., the line along which concentration of stress and local strain were most extreme. Crack initiation in backing steel occurred frequently at the toe of carbon steel weld metal. If the bottom of stainless weld metal were made smooth with cladding surface, the rupture time of transverse welded specimen would become identical to base metal.
The test results at 500°C, however, showed 45 degrees shear type fracture which is different from that at 600°C in which fracture occurred along bond.

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It is a well-known fact that ferritic steels are practically more useful than austenitic steels as tube materials, and it requires on the part of the users in general to raise as much as possible the creep resisting temperature of ferritic steels.
However, 21/4 Cr-1Mo steel has been so far the highest class of materials of ferritic steel in practical use, obtainable in market and used most frequently. On the other hand, still higher class materials, such as 7Cr-1Mo or 9Cr-1Mo steel, have been under researches and have been developed but they have not yet attained the stage for practical use.
In the super heater tubes for boiler with a steam temperature of 570°C, austenitic steels are used in general, because, when the metal temperatures of the tubes exceed 600°C, the strength of Cr-Mo steels should become insufficient. However, with the advancement of the development of ferritic steels, there is the possibility of either 7Cr-1Mo or 9Cr-1Mo steel coming into practical use up to the metal temperature of 570-620°C.
In the case of the temperature exceeding 600°C, 18-8 Ti steel used for supperheater tubes in America but some cases of accidents have been reported on account of the expansion and destruction of tubes while in use.
Causes of these accidents are due to such special properties of austenitic steels as difficulty to control grain growth and embrittlement of grain boundaries by precipitation of carbides. Those events could be prevented if there appear ferritic steels which can be used in practise in the temperature range. of 570-620°C.
In that sense it is desirable that ferritic steels with higher heat resistance be developed, and as the first step of such development, it is necessary to ascertain the maximum heat resisting temperatures of 21/4Cr-1Mo steel tubes, which are always determined according to the ASME or JIS Boiler Construction Code. However, these code values are estimated from the creep data of comparatively short-time and we understand that so far there have been no researches for determining the adaptability of code values to the steels on sale from the results of long-time creep data.
In the present studies, 10⁵hr 1% plastic strain stress and 10⁵hr creep rupture stress are estimated on the basis of 1-3×10⁴hr long-time creep data for 21/4Cr-1Mo steel, and from these results maximum allowable stress is determined and compared with code values.
The obtained results are given below:
1) Descending of creep strength in high temperature is much more rapid than that of code values.
2) At the temperature lower than about 570°C, creep strength is always over code values but is lower at the temperatures higher than that, and in the latter case it is necessary to select the maximum allowable stress lower than code values.
3) It is considered to be reasonable practically that heat resisting temperature of 21/4Cr-1Mo steel is 570°C.
It is also well-known that creep data are accumulated as Master curve by means of the method of Larson & Miller based upon the rate process theory, and the author advocates the possibility of the structural changes occurring during the holding of high temperature to be described on this Master curve by using the same parameter.
As an instance, the case of 21/4Cr-1Mo steel was taken up exmple. From the isothermal diagram obtained by annealing the steel within the temperature range of 400-750°Cfor 1000hr, the author sought the structural change diagram by using P=T(23+logt). And by imposing in upon the Master curve based upon the creep data of 1-3×10⁴hr, the tendency of strength changes can almost exactly be explained in relation to structural changes.
The author named this creep strength-structural change diagram the“Double Diagram”.

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We have practiced high temp. fatigue tests and creep rupture tests on seven kinds of heat resisting steels, and have examined the method of obtaining the allowable stress value and also various phenomenon incidental to these tests. Some of these specimens contains a chemical content non-standardized, but we have given them regular heat treatment in our tests. Fatigue tests of rotary-bending type was practiced at a revolving speed of 3000rpm, aiming at N=10⁷-2×10⁷ cycles, and creep rupture tests were carried out up to 1000-2000 hours.
The results of the above tests are tabulated in this report, and the high temp. strength of respective test specimens are found to be the values generally acknowledged so far and they involve nothing particular to explain. Here, we only wish to add our opinion concerning the method of assuming the allowable stress against long term of life and also a few points we noticed about it, which are as follows:
(1) In the high temp. fatigue, it seems that endurance limn of Ferritic Heat Resisting Steel is not observed above 400°C and that of Austenitic Heat Resisting Alloy above 600°C, and at higher temperatures than these degrees, it is considered preferable to obtain strength equivalent to 10⁵ hours by straight line extrapolating of S-N curves on the double logarithmic graph and to regard this strength as the allowable stress against long term fatigue.
(2) When we tried to extrapolate 10⁵ hours creep rupture strength from 1000 hours data, it was noted that about 30% difference of strength was recognized among direct extrapolation on the double logarithmic graph, Larson-Miller Method and Manson-Haferd Method. Here it is to be noticed that Manson-Haferd Method, the method which is reputed nowadays as the most accurate one, brought about the lowest value, it is deemed the safest in the process of obtaining the allowable stress to use Manson-Haferd Method, although careful examination of this method is indispensable.
(3) Besides, it is very convenient to draw the Fatigue-Creep Rupture Diagram in the form of a graph of allowable stress at high temperatures. If we plot on both axes of graph 80% of the 10⁵ hours value extrapolated by the aforesaid method and connect them with an elliptical curve, an almost proper graph could be made.
(4) Reduction of area at the creep rupture are apt to change regularly in accordance with the applied stress and the rupture time, To this also, we think the arrangement of master curve style may be applied. For trial, if we arrange the varying value in reduction of area according to the parameter of T(C+σ), we will be able to find a tendency that reduction of area at each temperature can mostly be gathered into a single curve. Further examinations would be necessary on this matter, but if they can be gathered into one single curve, it will be much convenient for estimating the variation of ductility for the long term.
(5) Further, it was shown that there are differences between high temperature fatigue and creep rupture in their precipitation and rupture conditions, and that change of hardness is so complicated, having relations not only with time of heating but applied stress and value of deformation.
The foregoing descriptions involve many problems and further examinations are required, and we expect the future study and discussion will be quite helpful to solve these ploblems.

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The authors have conducted some analyses on the dynamic creep for several years with the aim to estimate the dynamic creep strength from the informations on static creep, and verified the applicability of the analyses to several materials. In these studies, the analyses were confined to the case of the dynamic creep under an axial varying stress which was uniformly distributed over the cross-section of specimen. In the actual service conditions, however, there would be the cases of non-uniform stress distribution such as occurring in the members subjected to bending or torsion and also in the turbine blade, which is assumed to be under the combined stress state of axial static tension and alternating bending. Taking up the above-mentioned combined stress condition for turbine blade as a typical example of the cases of non-uniform stress distribution, the authors have carried out experiments as well as analyses, which brought the possibility of predicting the dynamic creep strength under this combined stress condition from static creep data. In this paper, the analyses are extended to the cases of bending dynamic creep, torsional dynamic creep and also of dynamic stress relaxation.
The outline of the analyses is as follows:
The assumptions which were used and verified in the previous study on the dynamic creep under combined static tension and alternating bending is utilized in this study also, that is,
(1) The equivalent static stress σ_e introduced for the purpose of predicting the axial dynamic creep strength in the transient stage of creep is also applicable to the second stage of creep.
(2) The distribution of alternating stress σ_a may be regarded as elastic, since the alternating component of strain is almost purely elastic as a result of the sufficiently high speed of stress alternation.
(3) The distribution of the equivalent, static stress σ_e under a dynamic stress condition is the same as that of the static stress which would produce the same creep.
If these assumptions are adequate the distribution of the dynamic stress which will produce the creep strain same as that under any static stress condition can be determined easily in the following way. The distribution of the equivalent static stress σ_e is obtained as a static creep problem according to the assumption (3). The distribution of the alternating stress σ_a is known as an elastic stress distribution from the assumption (2). On the other hand, the relation of σ_e, σ_a and σ_m is obtained from static creep data, and, finally, the distribution of the mean stress σ_m is determined from the relation, by inserting the above obtained value of the equivalent stress σ_e and the alternating stress σ_a.
According to the results of analyses on the bending dynamic creep and torsional dynamic creep, the relation between M_m/M_e and M_a/M_e in these cases (where M_m, M_a and M_e denote the mean moment, the alternating moment and the equivalent static moment, respectively, as in the case of stress) is similar. to the relation between σ_m/σ_e and σ_a/σ_e, provided that the latter relation for the case of the dynamic creep under an axial stress condition may be expressed approximately in a straight line within a sufficiently wide range of the variables.
In the case of stress relaxation, the differential equation for dynamic stress relaxation becomes as
1/Edσ_m/Idt+f(σ_m, ε₀-σ_m/E)=0,

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Compression-creep data for zircaloy 2, Mo-Cu-Zr alloy and 18-8 stainless steel were obtained at room temperature, 250°C, 316°C(600°F) and 450°C for a period of 100 hours. Zircaloy 2 and Mo-Cu-Zr alloy were casted respectively as ingot by the consumable-electrode double-arc melting. The test specimens were machined from a bar obtained from the ingot by forging, and annealed at 700°C (Zircaloy 2) and 750°C (Mo-Cu-Zr alloy) for 1 hour in vacuum furnace.
The test equipment for compression creep is the conventional tension creep machine with a fixture consisted of two yokes which convert tensile loading into compressive loading. The fixture used is of the similar type to the one developed at the Westinghouse Research Laboratory by M.J. Manjoine. The compression-specimen which has a diameter of 12mm and an overall length of 36mm was compressed between two seats, the ends of the specimen and of the seats being ground and lapped. The relative displacement of the yokes was measured by dial gauge extensometer as a measure of the strain in the specimen.
For checking the magnitude of instantaneous strain in creep tests, short-time tension and compression tests were made for zircaloy 2 and Mo-Cu-Zr alloy at 316°C by using the test equipment above-mentioned. The continuous loading was given by moving a running weight sliding on the loading lever arm of the creep machine.
Although at room temperature zircaloy 2 and Mo-Cu-Zr alloy have smaller instantaneous and creep strain in comparison with 18-8 stainless steel which displays appreciable creep at room temperature, they tend to have poorer creep resistance at higher temperatures, and the steady-state creep component becomes conspicuous for Mo-Cu-Zr alloy at 316°C and for zircaloy 2 at 450°C. The creep strength of Mo-Cu-Zr alloy at 450°C is stronger than that of zircaloy 2 when the stress level is below 17kg/mm².
Comparison of tension-creep and compression-creep properties for Mo-Cu-Zr alloy (at 316°C and 450°C) show that the alloy has poorer resistance in compression than in tension within a certain limit of stress, above which an effect of decrease of stress resulted from the increase of cross-section of a compression specimen would appear. The similar phenomena for S 816 and nimonic 90 at 1600°F have been reported by L.A. Yerkovich.
This difference in creep-resistance may partly be explained by the anomalous variation of the stress-strain relationship in tension and compression. But it should be taken into account as well that the bedding-down of the ends of compression-specimen and the anisotropic effects in the resistance to deformation produced in the process of preparing the test specimen are related to the difference in creep-resistance, although in our experiment the bedding-down of the compression specimen was minimized by lapping the ends of compression specimen.

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The hot torsion test has been adopted as an important means for the determination of a proper temperature for rotary piercing operation in the manufacture of seamless steel tube.
However, according to the means heretofore applied in which the workability of steel is judged by the number of twists to failure, a comparison of ductility at high temperatures is qualitatively possible, but it is almost impossible to judge the occurrence of crack in working from that number of twists to failure. Secondary stress caused by torsion test being considered one of the important factors for occurrence of cracks in the making of steel tube, the author has tried to verify the fact by making experimental studies on seven kinds of carbon steel and five kinds of stainless steel so that the result may be applied to the determination of hot workability, and has confirmed the following matters:
(1) Secondary stress which occur in torsion tests acts generally as tensile stress on the part of comparatively higher temperatures and as compressive stress on the part of lower temperatures. The temperature at the time of changing from tensile stress to compressive stress, that is, transition temperature varies with the kind of material.
(2) The value of secondary tensile stress differs according to testing temperatures. There is a summit on the correlative curve of these two factors, and the shape of this part varies with the kind of steel and its components. In practice, a higher summit and a wider temperature limit can be acquired in case of more easily workable materials.
(3) The temperatures where secondary tensile stress and the number of twists to failure respectively show the highest value do not necessarily coincide and defferent according to the kind of material.
(4) Secondary tensile stress varies with strain rate when another testing condition is constant, and at the same time, in some kinds of steel, the highest value, which is not clear in the measurement by the number of twists to failure, is shown by this strain rate. They are, therefore, important factors for the study of strain rate and cracks occurred in the process of rolling.
(5) Therefore, from the results mentioned in (1)-(4), secondary tensile stress has a close relation with ductility of materials. The secondary tensile stress is possibly applicable to the determination of cracks caused by working, which is difficult for us to judge by the number of twists to failure.

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Recently, AISI 321 type 18-8 Ti stainless steel has widely been used as the material of superheater or reheater tubes for large power plants, and it is generally recognized that the creep rupture strength of this type of steel varies greatly with the grain size in connection with heat treatment. However, in view of the fact that both of the grain size and condition of precipitation are affected by the heat treatment, we consider that the grain size is not a sole factor to have an effect on the creep rupture strength, but also the condition of carbides or σ phase precipitation has a great effect on it. To confirm this we carried out 650°C creep rupture tests of the material treated at various heat treatments after working. As a result of the tests, we have reached the following conclusions:
It makes a great difference in the creep rupture strength whether the heat treatment after working is carried out at a comparatively low temperature in the precipitation range, (less than about 1000°C), or at a high temperature in the solution range.
During recrystallization, precipitation and coagulation develop very quickly, and the precipitation hardening effect that prevents creep deformation is lost due to over-aging. Therefore, the strength of the material subjected only to working or treatment at a low temperature is small at the long period side. On the other hand, the strength of the material treated at a high temperature after working is not reduced, because the precipitation never occurs, and precipitated carbides or σ phase are dissolved in matrix during heat treatment.
There will be some relation between the creep rupture strength and the grain size. However, when the material is treated at a comparatively low temperature and has fine grains, it will also have an undue condition of precipitation. This coincidence might often cause a misconception that the creep rupture strength was affected only by the grain size, but we believe that the greater part of the reduction of the rupture strength is due to undue condition of precipitation. We shall make further studies on this point.

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Many researches on the high temperature strength of heat resisting alloys related to the microstructures have been presented in recent years. As a matter of fact, the correspondence of high temperature strength to the microstructures is very significant for precipitation type alloys. The heat resisting alloy A 286 reported in the present paper is one of the precipitation hardening alloy. As the first part of the studies on the alloy A 286, the aging structures are hereby presented.
As the results of phase identification by means of X-ray and electron diffraction, γ' phase, η phase, G phase, Laves phase, and TiC were identified. Nucleation and growth of these precipitates at a series of aging temperatures were observed. The results are as follows: G phase with the cubic type lattice was observed at the grain boundary, which did not significantly grow at 650°C, but remarkable growth occurred at temperatures above 760°C. η phase with the hexagonal type lattice nucleated at the grain boundary and grew into the grain with Widmannstätten appearance, which was not observed at 650°C aging, although appreciable η phase precipitation was observed at 718°C for 50hrs and at 800°C for 2hrs, respectively. γ' phase with the f.c.c. type lattice was observed within the grain with the spherical shape, of which many of small precipitates were nucleated at 650°C, although its rate of growth was remarkably increased as the temperature become higher. It might be considered that the hardness was mainly determined by the precipitates of γ' phase. Laves phase with the hexagonal type lattice was observed at the grain boundary and within the grain, which precipitated at the stage of over aging. TiC was stabilized at the solution treatment, and subsequently it was not essentially changed during the aging treatment.
Effect of the melting processes on the precipitaion process was significant. Especially, the effect was very remarkable at the time of precipitation of G and η phase at the grain boundary in the first stage of the aging, that is, the rate of precipitation in vacuum melted alloy was smaller than that in the air melted one. As for the intragranullar precipitates such as γ' phase and Laves phase, the same effect was also recognized, although the order of magnitude was relatively smaller. It is supposed that the phenomena as mentioned above is due to the existence of boron atoms which came into the solution from the crucible during the melting process.

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The present study was carried out to clarify the high temperature properties of stainless steel welded joints of AISI types 347 and 321 in as-welded condition, and the effect of post-heating on the creep rupture properties of these joints.
To prepare test pieces, stainless steel plates of types 347 and 321 with 25mm thick were welded using 16 Cr-8Ni-2Mo and 19 Cr-8Ni-1Mo electrode, Some of these welded joints were given post-heating at 1050°C for 1hr and at 900°C for 2hrs.
In the tension test, specimens as-welded fractured at base metal, and their tensile strength was equal to that of base metal. Tensile strength of specimen post-heated was slightly lower than when it was left as-weld. (Fig. 4, 5)
Type 321 stainless steel welded joint using 16-8-2 electrode had the best fatigue properties among the joints in as-welded condition. (Table 2)
In creep rupture test, type 347 stainless steel welded joint as-welded fractured at the heat-affected zone of base metal and its rupture strength was very poor. However type 347 stainless steel welded, joint recovered considerable rupture strength by post-heating at 900°C and 1050°C. (Table 4) It seems that this decay of the heat-affected zone is caused by the strain, near the weld groove. By post-heating at 1050°C, coarsening of grain and softening occurred at the heat-affected zone. (Fig. 3, Photo. 1, 2) It seems that this zone has been slightly hot-cold worked during weld by thermal stress caused from temperature gradient and contraction of weld metal. On the other hand, type 321 stainless steel welded joint as-welded fractured at base metal and showed no decay such as shown by type 347 stainless steel welded joint in the rupture test. 1050°C post-heating make its weld boundary so weak, that the creep specimen ruptured at the boundary along the bond. But, 900°C post-heating didn't produce such a decay. (Table 4, Photo. 4, 5) It is not clear why such a decay occur in type 321 stainless steel welded joint.

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The strength in thermal stress cycling is one of the most complicated in the strength of material at high temperature, as the authors pointed out in the previous report. The study on such strength is important from fundamental and also practical point of view. Practically the structure members subjected to load at elevated temperature are not in a steady condition but in unsteady condition, as for stress and temperature, especially, variation of thermal stress is the most severe in the unsteady conditions. On the other hand, in recent years, attention has been placed on such an aspect of the fatigue problem that failure results from comparatively small cycles of severe stresses where strain amplitude of every cycle is the independent variable which destines the life. Also, attention has been given to the problem of influence of temperature variation on fatigue life. These attentions are focused on the problem of thermal stress cycling.
The authors have been interested in the problem regarding the meaning of strain amplitude in thermal cycling test and in the study of the relationship between strength in thermal cycling and the fundamental strength at elevated temperature, such as static tensile strength, strength in strain-cycling, strength of creep rupture and others. The study on such relationship is practically important to make predictions as to test data under thermal cycling from the informations of other fundamental tests.
The test of thermal cycling descrived above is that of completely reversed thermal stress cycling. However, the test of pulsating tensile thermal stress cycling is also necessary in order to know the relationship between the strength in thermal cycling and that of creep rupture. This is necessary to obtain the information of strength of material subjected to both mechanical steady stress and cyclic thermal stresses, from a practical point of view. Therefore, in the present paper, a test of pulsating thermal stress cycling on AISI 304 type stainless steels is performed, using the same testing apparatus as in the case of conventional thermal cycling test, being remodelled in gripping device so as to apply specimen tensile thermal stress only, free from the compressive stress. The test results under pulsating tensile thermal stress cycling are discussed in comparison with those under fully reversed thermal stress cycling. Attention is also placed on the relationship between the strength in pulstaing tensile thermal stress cycling and that of static creep rupture.
Following are conclusions from the present study:
(1) Failure life for pulsating tensile thermal stress cycling is much smaller than in the case of fully reversed thermal stress cycling, that is, the failure life under completely reversed cycling is nearly two hundred times as much as that under pulsating cycling as for the basis of total strain amplitude.
(2) The life for pulsating tensile thermal stress cycling is a little smaller than that estimated from the criterion regarding life of creep rupture of the material subjected to simultaneous periodic variation of temperature and applied stress (Eq. 4), that is, the estimated result is nearly three times as big as the experimental one. It would mainly be because of neglecting exact calculation on the increase of thermal stress amplitude expected in accordance with the number of cycles in the criterion. Thus, in such a case of prediction, the effect of increase of thermal stress amplitude on the rupture life must be taken into consideration

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We have investigated on the changes in dimension, hardness, ultrasonic attenuation and microstructure of three kinds of carbon steel (E, D and J) having different amounts of non-metallic inclusion MnS during simple thermal cycling between 720°C and 50°C, and further, carried out the same studies on armco iron.
The results were as follows:
1) The sample E which contains a small amount of inclusion elongated in the longitudinal direction, but the samples D and J with higher quantities of the inclusion contracted after thermal cycling. On the other hand, their diameters were found to expand regardless of inclusion quantities. Therefore, some decrease in specific gravity has taken place in all samples.
2) The changes in hardness, grain size and ultrasonic attenuation mostly occurred during the period of 50 cycles.
3) The changes in hardness, grain size and ultrasonic attenuation in all samples which were subjected to isothermal annealing for 310hrs at 720°C have the same tendency with the case of thermal cycling, but the dimensional change could not be recognized.
4) The dimensional change in armco iron was small compared with the sample E, but plastic deformation and formation of sub-grain due to thermal stress was clearly observed.
The anomalous changes in dimension and mechanical properties are supposed to be caused by reason of that the concentration of the thermal stress at the boundary regions between grain and inclusion induced by heating and cooling has brought about some plastic deformation, which results to the origin of crack or various defects. The behavior of materials bearing inclusion are influenced by the form, properties, size and distribution of the inclusion. Furthermore, the chemical and physical reaction between base metal and inclusion is to be accelerated by thermal stress, and consequently, various anomalous phenomena might be appeared.

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We obtained the following results from the short-time tensile test, creep-rupture test and fatigue test on the high temperature properties of the commercially pure titanium and zirconium:
(1) Commercially pure titanium has a comparatively higher strength considering its small specific gravity, and a comparatively higher yield strength considering its tensile strength.
(2) Commercially pure titanium has a unique creep property in the range of 200°C to 300°C with the existence of limit stress in creep test at the said temperature range. The same tendency occurs in commercially pure zirconium, but it is not so remarkable as in titanium.
(3) The fatigue limit of commercially pure titanium in the high temperature fatigue test tends to decline sharply in the temperature range of room temperature up to 200°C and then gradually in the temperature range of 200°C to 400°C by temperature elevation, as the decline tendency by temperature elevation of the short-time tensile strength.
(4) The unique creep property of commercially pure titanium in the temperature range of 200°C to 300°C is proved to be the result from strain aging, which causes the decline by temperature elevation of the short-time tensile strength to become slow in the temperature range of 250°C to 500°C.

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The physical and mechanical properties of molybdenum are superior to those of the other metals at high temperature. That is, the melting point of molybdenum is very high (2620°C) and accordingly the recrystallization temperature becomes high (approx. 1200°C). Also, the Young's modulus is high, and even at 870°C, the value is 1.3 times (28100kg/mm²) as high as that of a carbon steel at room temperature. The coefficient of thermal expansion is small (5.5× 10^-6/°C) and is about a quarter of that of Fe-Cr-Ni alloy. The creep strength is about 17.5kg/mm² at 980°C and is higher than those of the conventional heat-resistance alloys. The tensile strength and yield point decrease little when the temperature rises.
As described above, molybdenum has many superior properties as the heat-resistance materials, but it is a big defect that the oxidation resistance at high temperature is much inferior. That is, the oxidation of molybdenum proceeds, up to 490°C, according to the parabollic law, but above this temperature, the vaporization of the latest product (MoO₃) occurs, and oxidation of molybdenum proceeds all the more rapidly so as to provide the vaporization products. This phenomenon becomes remarkable with the increase of temperature.
Accordingly, when molybdenum is used as a heat-resistance material, it is necessary to let the oxidation resistant metal elements alloy with molybdenum or to perform the oxidation resistant coating on molybdenum. Concerning the former method, it is necessary to add a large amount of alloying elements to molybdenum, and consequently the mechanical properties of molybdenum at high temperature are affected. Therefore, the latter method is desirable, and many investigations have been made hitherto.
We investigated the high temperature oxidation protection for molybdenum by means of siliconizing, chromizing or Ni, Cr and Co electro-deposit coating, etc. and found that Cr and Ni double layer deposit coating was the most adequate among the oxidation protective coatings for molybdenum, but the oxidation resistance life of molybdenum at 1000°C in air obtained by this coating was up to 300hrs.
In this report, through studies on the cause of destruction of this coating at high temperature oxidation, a considerable effort was made to obtain Cr-Ni coating with a longer oxidation resistance life, by improving the electro-deposition process, that is, anodic treatment for molybdenum surface before chromium deposition, and expelling absorbed hydrogen in Cr-deposited layer that were found to be the cause of blistering and spalling of Ni layer.
By this improvement, the double deposited layer showing oxidation resistance life, at maximum, 700hrs at 1000°C in air was obtained when Cr layer was 0.025mm and that of Ni 0.150mm thick.
The mechanism of oxidation of this coating was also investigated. The oxidation resistance of this coating mainly owes to the outer Ni layer, while the inner Cr layer has nothing directly to do with the oxidation resistance. The rate of weight gain of the specimen with double coating by oxidation is slower than that of pure nickel because of the diffusion of chromium into nickel layer by heating. It was found by X-ray diffraction and qualitative analysis that the oxidation product on coated surface was not nickel oxide (NiO) but solid solution with face centered cubic lattice consisting of Cr, Ni and O.

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