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By Sir Robert Hadfield

LAST year this Institute was good enough to accept some remarks by the writer regarding sound steel, entitled Plant for Hadfield Method of Producing Sound Steel Ingots, being a continuation of a research in which for some time the writer has been engaged, and full particulars of which were submitted in papers entitled A New Method of Revealing Segregation in Steel Ingots, and Method of Producing Sound Ingots, read before the Iron and Steel Institute in October, 1912. These papers gave rise to a very interesting discussion, both in England and America. This question is one of vital importance, as unless sound steel in the form of ingots is first obtained, sound steel in the finished forms of rail billets, rails, bars, sheets, and other forms cannot be expected. The working stresses demanded by modern conditions are so great, each year they probably approach nearer to the limit of safety, that without doubt those who produce steel will be compelled to pay the attention which this subject deserves. Moreover, the problem is not an insuperable one, as shown by the writer in the papers above referred to. In view of the information which has already been presented in these papers, it is not now necessary to go into details of the system advocated by the writer, except to repeat that the ingots and steel now referred to were produced by employing in the first place sound fluid steel, free from blowholes, this being poured into ordinary ingot molds, provided with the necessary feeding heads, the insulating slag medium, also charcoal and air blast or other suitable form of applying heat to the head portion of the ingot. Tile upper or head portion of the ingot is kept fluid by the intense heat generated by the air blast and charcoal, so that such fluid steel passes into the body of the ingot below the head, thus feeding and avoiding the shrinkage or piping which would otherwise occur. The result is the production of a sound steel ingot, free from blowholes, segregation and piping, and in which the waste portion, usually termed the discard, is but little over 7 or 8 per cent.

Jan 4, 1914

By Charles S. Barrett

Recrystallization textures—thc orientations of grains after recrys-tallization—have been studied extensively not only because of their metallurgical importance but also because of the information they yield regarding the atomic movements during reerystallization. Recrystallized grains must possess the same orientations as the submicroscopic nuclei from which they have grown, and thus they offer a direet road by which experimenters can penetrate the submicroscopic realm. Many perplexing results have come of these studies, but one of the most striking has been the apparent difference in behavior between single crystals and polyerystalline grains. For example, Burgers and Louwerse have shown1 that when single crystals of aluminum are deformed by compression and are reerystallized (at 600' C.) the new grains appear with orientations that are different from the orientations of the deformed crystal in which they grew; but polycrystalline specimens, on the other hand, appear to have a [110] fiber texture both before and after reerystallization and thus to retain their texture during recrystallization. This apparent difference between the habits of single crystals and small grains is the subject of the first part of the present research and has been explained by findings reported herein. A similar situation appears to exist for aluminum when elongated or drawn into wires, for single crystals alter their orientation 2,3 and polycrystalline wires retain their texture"5*" upon recrystallization after elongation (at least with high-purity aluminum), and it is possible that the explanation found in this paper for compression can be extended to cover elongatioil and other types of deformation as well. The second part of this paper deals with the theory of reerystallization textures. In the most successful theory for aluminum reerystallized after compression,1 much weight is given to the orientation of highly stressed nuclei in the deformed material. Taylor's "local distortions" arc assumed to exist along the slip planes where displacement has occurred, consisting of fragments rotated in a specific way with reference to the slip plane and slip direction; namely, about an axis lying in the slip plane perpendicular to the slip direction.1,7,89

Jan 1, 1940

By J. W. Spretnak, G. W. Cunningham

A study has been made of the growth and disappearance of voids in the interface of pressure bonded copper. The formation of visible voids is primarily due to spheroidization and coalesence of microvoids produced during the application of pressure. The microvoids could be prevented from forming by the application of high pressures or by increasing the ductility of the copper. Migrating grain boundaries were found to be more effective in void removal, than stationary grain boundaries. DURING the course of an investigation on the mechanism of pressure bonding, it was found that voids formed in the bonding interface which inhibited migration of the interfacial boundary and grain boundaries in the bond region. Therefore, an effort was made to determine the nature of the voids and how they form, grow, and disappear. This paper reports the results of such an investigation. The existence of interfacial voids is not solely a characteristic of massive components of metal which have been bonded by the application of heat and pressure, but is a phenomenon which has been observed in sintering as well. Simonsen attempted to study the sintering process by hot-stage microscopy and thereupon observed the formation of interparticle porosity as sintering proceeded. It was assumed that the porosity was due to grooves created at the location where grain boundaries joined the surface of the particle. Pranatis, Castelman, and seigle2 have also observed voids along the areas of contact in wound-wire specimens used for sintering studies. They did not discuss the point in detail since this behavior was not general, but they did suggest that the voids might develop as a result of stresses along the interparticle contact area. In previous investigations on pressure bonding, workers3 have suggested that voids may exist in the region where mating surfaces are placed in intimate contact, but the existence of the voids was not verified and no attempt was made to study the growth and disappearance of these voids. Durst noted the existence of pores in roll-clad structures prepared by cold reducing two sheets to force them in contact and then annealing to improve the bond. Since his unbonded regions contained a gas which he assumed to be insoluble, he proposed a mechanism which consisted of spheroidization of the unbonded regions and a growth of the contact area, but no speculations were made regarding disappearance of the gas filled pores. One interesting comment made by Durst was the suggestion that flattened voids with sharp corners which are produced during roll bonding will have a greater driving force for spheroidization than the more rounded pores created by pressing metal powders and thus spheroidization and neck growth will occur at lower temperatures. In this study it was necessary to demonstrate conclusively that voids can actually exist in the interfacial bond area. Experimental work was then directed toward a study of the mechanism whereby the voids grow and subsequently disappear.

Jan 1, 1963

By C. H. Samans, G. F. Tisinai, J. K. Stanley

The addition of nitrogen (0.10 to 0.20 pct) to Fe-Cr-Ni alloys of simulated commercial purity results in a real displacement of the u phase boundaries to higher chromium contents. The effect is small for the (Y + s)? boundary, but is pronounced for the (y + a +s)/(y + a) boundary. Although there is an indication of an exceptionally large shift of the n boundaries to higher chromium contents, especially in steels with nitrogen over 0.2 pct, the major portion of this apparent shift results from the fact that carbide and nitride precipitation cause "chromium impoverishment" of the matrices. The effect of combined additions of nitrogen and silicon to the Fe-Cr-Ni phase diagram is demonstrated also. Nitrogen can nullify the effect of about 1 pct Si in shifting the (y + o)/? phase boundary to lower values of chromium at all nickel levels from 8 to 20 pct. NItrogen can nullify this U-forming effect of about 2 pct Si at the 8 pct Ni level, but not at the 20 pct Ni level. The alloys studied were in both the cast and the wrought conditions. There are indications that the u phase forms more slowly in the cast alloys than in the wrought alloys if both are in the completely austenitic state. The presence of 6 ferrite in the cast alloys accelerates the formation of U. Cold working increases the rate of o formation in both cast and wrought alloys. THE major improvement in Fe-Cr-Ni austenitic alloys in recent years has been in the addition or removal of minor alloying elements to facilitate better control of corrosion resistance, sensitization, and heat resistance. One shortcoming of the austenitic Fe-Cr-Ni alloys, which never has been completely circumvented, is their propensity toward u formation. In the AISI-type 310 (25 pct Cr-20 pct Ni) and type 309 (25 pct Cr-12 pct Ni) steels, sufficient amounts of u phase can form, if service or treatment is in a suitable temperature range, to cause severe embrittlement. Also, there is a growing conviction that this phase may be contributory to some unexpected decreases in the corrosion resistance of certain of the 18 pct Cr-8 pct Ni-type steels. The present paper discusses the effect of nitrogen additions on the location of the (r+u)/d and the (y+a+u)/(y+a) phase boundaries in the ternary Fe-Cr-Ni system, for cast and wrought alloys of simulated commercial purity, and in similar alloys containing up to about 2.5 pct Si. The objective is to define compositional limits for alloys which will not be susceptible to u formation when used near 1200°F (650°C). An excellent review of the early studies of the u phase in the Fe-Cr-Ni system has been compiled by Foley.1 Rees, Burns, and Cook2 have determined a high purity phase diagram for the ternary system, whereas Nicholson, Samans, and Shortsleeve3 are- stricted themselves to a portion of the simulated commercial-purity phase diagram. Both groups of investigators show almost an identical position for the commercially significant (y+u)/y phase boundary. Further comparison of the work of the two groups indicates that, below the 8 pct Ni level, the commercial alloys have a decidedly greater propensity toward u formation than the high purity alloys. The two groups of workers agreed that both the AISI-type 310 (25 pct Cr-20 pct Ni) and the type 309 (25 pct Cr-12 pct Ni) steels are well within the (y+~) region and that the 18 pct Cr-8 pct Ni-type alloys straddle the U-forming phase boundaries. Nicholson et al.3 showed, in addition, that these boundaries shift toward lower chromium contents if greater than nominal amounts of silicon or molybdenum are added. The effect of nitrogen on the location of the s phase boundaries in the Fe-Cr-Ni system has not been known with any certainty. In 1942, an approach to this problem was made by Krainer and Leoville-Nowak,' but at that time they apparently were unaware of the slow rate of s formation in strain-free samples and aged their samples for insufficient times, e.g., 100 hr at 650°C (1200°F) and 800°C (1470°F). For this reason, it would be expected that their (y+ u) /y boundary would be shifted toward lower chromium contents (restricted ?-field) when equilibrium conditions were approximated more closely. Procedure for Studying the Alloys The alloys used were prepared in the following way: Heats of 200 lb each were melted in an induction furnace. A 5 lb portion of each heat was poured into a ladle containing an aluminum slug for de-

Jan 1, 1955

By R. Maddin, A. Lawley

Single crystal Mo-Re alloys (99.99+ purity), grown by electron bombardment floating zone heating, were deformed in tension at temperatures from -196° to +200°C. At -196oC, Mo-6 pet Re and Mo-20 pct Re alloys deform by slip on systems of the type (101) [1ll]. Above-78oC, the operative slip system is of the type (112) [111], (101) [1ll], or (211) [1ll], depending upon orientation. Mo-39 pct Re alloys deform by twinning and slip at all test temperatures with a (112) [111] type twinning system. Critical resolved shear stress for slip and twinning is calculated as a function of composition and temperature, and the mechanical properties compared with those of less Pure polycrystalline alloys. HEXAGONAL close-packed rhenium, c/a = 1.615, is highly soluble in molybdenum. The bee molybdenum-rich solid solution extends to approximately 42 at. pet Re* at 2500°C,' Fig. 1. Recent work on polycrystalline alloys has shown conclusively that the addition of rhenium to molybdenum improves both the fabricability and the low-temperature ductility.2"5 Similar behavior is observed for W-Re and Cr-Re alloys.5 Initially, Geach and Hughes2 found that the Mo-30 pct Re alloy could be cold rolled to approximately 90 pct reduction before the onset of cracking. The remarkable effect of the rhenium was later confirmed by Jaffee et al. 3-5 with an optimum condition of maximum strength combined with maximum ductility at the Mo-35 pet composition. At higher rhenium contents, the ductility was severely impaired due apparently to the presence of brittle tetragonal s phase, Fig. 1. In the Mo-35 pet Re alloy, at least 110 ppm (parts per million) oxygen could be tolerated before cold workability was reduced. By comparison, the corresponding tolerance figure in unalloyed molybdenum is ~ 10 ppm. The effect of rhenium was considered in terms of: 1) a reduction in oxygen solubility in the lattice; 2) a redistribution of intergranular oxide from grain boundary films to isolated spheres; and 3) the promotion of twinning as a further mode of plastic deformation. Paradoxically, twinning is often associated with brittle fracture in iron and the silicon irons.6'7 Work to date has been restricted to are-melted and powder metallurgy polycrystalline material of ~ 99.9 pet purity. The present investigation was initiated for the purpose of determining the mode and crystallography of plastic deformation of a series of bee Mo-Re alloy single crystals over a wide temperature range from -196o to +200°C. Using high-purity zone-melted material (~ 99.99 + pet purity), the specific role of rhenium in the molybdenum lattice may be understood more fully since effects due to grain boundaries and oxide redistribution will be absent. At the same time, the study provides a comparison of single and polycrystalline deformation behavior. EXPERIMENTAL TECHNIQUES A) Preparation and Purity of Single Crystals. Alloys in the form of either are-cast and swaged or sintered and ground bar, 0.120 in. diam, were obtained from the Chase Brass and Copper Co. Alloy single crystals ~ 8-in. long were grown from the

Jan 1, 1962

By B. L. Averbach, M. Cohen, C. H. Shih

The isothermal formation of martensite is studied in Fe-Ni-Mn and Fe-Mn-C alloys under conditions where the athermal transformation is completely avoided, there being no martensite present at the beginning of the isothermal reaction. The initial rate of martensite formation is extremely low under these conditions, but it increases greatly as soon as some martensite is formed. It appears that true initial characteristics of the isothermal reaction can be studied only in samples wherein the formation of martensite on cooling is completely eliminated. THE phenomenon of isothermal martensite formation is well established1 -" but the observations have been made in the presence of varying quantities of athermal martensite. Complete avoidance of athermal martensite and the formation of martensite under pure isothermal conditions have been reported by Kurdjumov and Maksimova2 in an Fe-Ni-Mn alloy (23 pct Ni, 3.4 pct Mn, and the balance Fe) and have been confirmed by recent experiments of Cech and Hollomon in a similar alloy. In each case, the isothermal transformation rate was a maximum at the beginning of the transformation and decreased as the reaction proceeded. These initial transformation rates were used by Fisher7 for the confirmation of calculated nucleation rates based on classical nucleation theory extended to include the effect of elastic strain energy. Cech and Hollomon have admitted, however, that up to 3 pct martensite formed on cooling in their samples and thus was present at the initiation of their isothermal runs. In this paper, experiments on similar alloys are described, except that care was taken to remove the incidental martensite developed at the surface prior to the isothermal reaction. It is shown here that the reaction rate at the beginning of the strictly isothermal reaction is very low. As soon as some detectable amount of martensite is produced, the transformation rate increases rapidly but ultimately decreases as the reaction proceeds. When martensite is allowed to form isothermally in the presence of about 2 pct of prior martensite, the apparent initial reaction rate is a maximum and corresponds to the observations of Kurdjumov and Maksimova and Cech and Hollomon. It is evident that the cooperative nature of the martensitic reaction results in a significant stimulating or autocatalytic effect on the isothermal reaction rate. Experimental Procedure The chemical compositions of the alloys used in this investigation are listed in Table I. The Fe-Ni-Mn alloys A and B were melted in an induction furnace, cast into ingots 2 in. diam by 4 in. long, and forged into % in. diam rods. The manganese steel (alloy C) was melted in an arc furnace and cast into ceramic molds 1/4 in. diam by 5 in. long. The rods were sealed in evacuated Vycor tubes and homogenized at 1175" to 1200°C for 50 hr. The long homogenization resulted in some manganese loss from the surface, producing a thin layer of transformed material. This surface layer was removed by grinding and the specimens were then swaged to 0.050 in. diam wire. Measurements of electrical resistance were made with a Kelvin double bridge to follow the course of the martensitic transformation. A small permanent magnet was also useful in scanning an entire specimen for small localized quantities of martensite. This test is almost as sensitive as the electrical resistance method in detecting the presence of martensite (which is ferromagnetic in these alloys) and it was especially useful in deciding whether a specimen was completely free of martensite. The wire specimens (0.050 in. diam by 2.5 in. long) were sealed in evacuated tubes and austenitized for 5 min at 1150°C. The specimens were quenched into water and, unless otherwise stated, stress-relieved for 1 hr at 650 °C in order to induce greater uniformity in the subsequent martensitic transformation. As dem-

Jan 1, 1956

By E. I. Salkovitz, F. W. von Batchelder

DURING the last few years several papers1-' have been published in which internal friction measurements have been used to determine the quantity of carbon or nitrogen dissolved in a iron. This method is based on the fact that, in pure Fe-N and Fe-C alloys, nitrogen and carbon cause peaks in the internal friction vs. temperature curve; at a frequency of 1 cycle per sec (cps) these peaks appear at $-200 and +360C, respectively, and if internal friction is expressed as Q-1, which is equal to times the logarithmic decrement, this figure will, by coincidence, give the weight percentage of nitrogen or carbon in solid solution.' In most of the studies published, very pure iron (Puron) was used. In one case the influence of manganese on the behavior of nitrogen in a iron was mentioned.4 It would be of great value, especially for the study of aging and related phenomena, if this method could be applied to commercial mild steel with norrnal percentages of manganese, phosphorus, sulphur, etc. As a rule, in such steels carbon and nitrogen are present together, and the amounts of these elements in solid solution should preferably be determined separately. The fact that their internal friction peaks appear at different temperatures (+36° and + 20°C in pure iron) seems promising for such a selective determination. It is necessary, however, that the peaks are narrow enough not to interfere with one another, and further, that there is no effect from carbon on the behavior of nitrogen in stress induced diffusion, and vice versa. According to Fast,' however, in the presence of manganese the nitrogen peak gets wider and appears at a higher temperature nearer the carbon peak, and thus the prospects for selective determination are reduced. (The carbon peak was not affected by the presence of manganese.) But even if fully quantitative de- terminations of the two elements are not attained, some relative figures on the sum of them and an indication of which one is the predominant one, would be of value. Some results from experiments along these lines will be reported here. An internal friction measuring device was assembled mainly following a description by KG." The measurements were also carried out as reported in that paper. Table I gives the chemical compositions of the four different materials used in this investigation. All materials were drawn to wires with diameter about 0.8 mm, and the analyses refer to these final shapes. Determination of Nitrogen Peak Wires from steels A, B, and C were decarburized in wet hydrogen (about 30 pct H2O) at 700°C until no carbon or nitrogen could be found by chemical analysis. No peak was then found in the internal friction. The wires were further annealed at 580°C for 16 hr in hydrogen bubbled through a solution of 6.8 pct NH3 in distilled water. Chemical analyses gave 0.04 pct N in A and C, and 0.06 pct in B. After annealing at 580°C for 15 min in a lead bath and water quenching, the internal friction was immediately measured. The resulting curves are shown in Fig. 1. At these high values of the internal friction the accuracy is comparatively poor; at Q-1 - 0.030 it is estimated to be 0.002. In order to make possible a calculation of the activation energy for diffusion of nitrogen, measurements were made at two different frequencies. Steel A shows rather broad peaks, and their temperatures can hardly be expected to be very accurate. Steels B and especially C are better in this respect. The finding of Fast' that manganese broadens the nitrogen peak is con-

Jan 1, 1953

By W. A. White

UNDERCLAY squeeze is the plastic flowing of underclay below coal pillars into mined-out entries and rooms. Squeezes may be caused either by wet mine conditions where the moisture is taken up by the clay—making it more plastic—or by the presence of certain clay minerals, such as mont-morillonite, which are plastic when there is enough natural moisture. Under either condition, sufficient stress will cause the underclay to flow. Squeezes may fill an entry or room within several hours, or it may be several months before movement stops. Squeeze areas commonly cover hundreds or thousands of square yards. The Coal Div. of the Illinois State Geological Survey has collected samples of underclays from various mines and from cores drilled by coal companies in the state. From these samples natural moisture content, particle size and particle size distribution, and clay mineralogy are determined. Data from samples collected from the underclay below Herrin (No. 6) coal in the Lumaghi Coal Co. mine near Collinsville, Ill., are used in this report to clarify points in the discussion. Data from samples taken from other mines lead to the same conclusions. A generalized geological section for the area is given in Table I. The underclay below coal No. 6 ranges from a few inches to 3 or 4 ft thick. The upper 6 to 18 in. are noncalcareous; the lower part becomes more calcareous with depth. Limestone nodules become larger with depth, almost grading into the limestone below. In areas where the clay is nearly absent the limestone may thicken to take its place; thus they may complement each other in thickness. There are two types of squeeze areas in a mine. In one type up to 3 or 4 ft of underclay lie in what appears to be a depression in the limestone underneath. This is suggested by a rim surrounding the squeeze area where the underclay is thin and the limestone lies within a few inches of the coal. In the second type the underclay of the squeeze area is surrounded by underclay of similar thicknesses which does not squeeze. Test Methods: Natural moisture content, particle size and particle size distribution, and clay mineralogy were determined by the following methods: Natural moisture is the moisture content a clay contains in its natural environment, expressed in percentage of the weight of ovendried clay (110°C or 230°F). To obtain the natural moisture content, given in Table 11, the clay samples were sealed in containers as soon as the samples were taken in the mines. Particle size and particle size distribution were determined by the pipette method,' and the data were plotted as cumulative curves, Fig. 1, on semi-logarithmic paper. Particle size was plotted on the logarithmic scale and percentage on the arithmetic scale. The quartiles (diameters that correspond to frequencies of 25 and 75 pct) and medians (diameters that correspond to the frequency of 50 pct) of the samples were determined from these curves. The clay minerals were identified from the data obtained by differential thermal' and X-ray powder" techniques. The differential thermal analyses were made on the whole samples, whereas the X-ray analyses were made on fractions of less than 2p. Data and Discussion Moisture Content: After the sealed jars had stood for 24 hr, those that contained squeeze samples had perspired, so that droplets of water had accumulated on the insides of the jars. Table II shows the moisture in the clay after it had perspired for 24 hr. (This perspiring was not noted in samples from nonsqueeze areas.) After perspiration one of the squeeze samples still had a higher moisture content than the samples that did not squeeze. Clay Mineralogy: The thermal analysis, Figs. 2 and 3, indicates that the clay minerals are chiefly micaceous in type; the low-temperature endothermal

Jan 1, 1957

By R. M. Rose, D. P. Ferriss, J. Wulff

Single crystals of tantalum, tungsten, and the binary alloys thereof were grown by electron-beam zone melting and tested in tension between 77° and 373°K. The ductile-to-brittle transition temperature increased sharply with solute concentration, for both tungsten-rich and tantalum-rich alloys; apparently, the maximum transition temperature exists near the 50 pet concentration. Twinning was observed to be a consequence of cleavage-crack propagation. Cleavage occurred on(100)planes, and extensive tearing and "river markings" were present. The onset of brittleness could not be associated with any change in deformation mode, the yield point phenomenon, or the large rise in yield stress with decreasing temperature. Cleavage was observed to initiate at either the specimen surface, or subgrain boundaries. The observation was made that the grain boundaries of Polycrystalline materials are probably even better as nucleation sites for fracture, and are responsible for brittleness at high temperatures and lower solute concentrations than those of the corresponding single crystals. TUNGSTEN and tantalum, the two high-melting neighbors in the third transition period of the periodic table, present striking contrasts. Although both are bee, tungsten has far higher elastic moduli, and the highest ductile-to-brittle transition (DBT) of all the pure metals.' Tantalum, however, is ductile at liquid helium temperatures. Large strain aging and yield point phenomena are observed in tantalum,2,3 but only to a lesser extent in tungsten.4 Tungsten work-hardens at a high rate,5 whereas this effect is negligible in tantalum.' Since the mutual solid solubility of these two metals is complete,7 and no ordering has yet been reported in the binary system, the possibility of depressing tungsten's DBT by tantalum additions was considered. Also, tungsten additions to tantalum could produce hard, high-melting alloys which might be ductile at room temperature. Such alloys offered excellent bases for complex alloy development, particularly with regard to the effect of rhenium on Group VI metals.8"10 In addition, there was the opportunity to examine flow and fracture modes in the vicinity of the DBT, with particular emphasis on the correlation between brittleness and fundamental deformation mechanisms. Myers11 has reported that recrystallized alloys with up to 10 pet W in tantalum are ductile at room temperature, but alloys with 20 pct W could not be swaged at temperatures as high as 500°C. The whole spectrum of Ta-W alloys has been surveyed by Braun, Sedlatschek, and Kiefferl2 using sintered powder compacts. In addition to some physical properties they reported a maximum hardness at about 50 at. pet W, dropping to much lower hardness levels for the pure components. They were able to cold work tantalum-base alloys with as much as 30 pet W and to warm work tungsten-base alloys with up to 20 pet Ta by procedures similar to those used for pure tungsten. EXPERIMENTAL In order to study deformation modes and to keep the impurity concentrations at a minimum, electron-beam zone melting was used to prepare single crystals of high purity. The electron-beam zone melting apparatus is similar to that described by Calverley, Davis, and ever" and incorporates a grid control rectifier to maintain control of emission current. Pure metals were refined by making three zone passes at 4 mm per min up the rod. Alloys were prepared by melting together wire of one constituent on a previously melted rod of the major constituent. Subsequently, two zone passes were made in opposite directions to level the composition after the manner suggested by pfann.14 Purification of the starting materials was accomplished by evaporation of the several impurities, and the degree to which this was accomplished is indicated in the chemical analyses of Table I. The uniformity of composition along a zone levelled rod was within the i 1 pet accuracy of the fluorescent X-ray technique used to determine composition. Crystallographic orientation of the

Jan 1, 1962

By J. P. Hammond, C. J. McHargue, S. E. Adair

Cold-rolled sheet textures have been determined for binary solid solutions of aluminum, columbium, tantalum, and zirconium in titanium. An alloy containing 3.8 pct Al had a (0002) [10101 texture, whereas alloys with the other elements had rotated (0002) [10101 textures similar to that of iodide titanium. THERE have been relatively few studies of the effect of alloying on deformation textures. A second element may be present in the base metal in solid solution and/or as a second phase. The literature offers no satisfactory explanation for the effects of solid solution alloying on textures noted thus far. For example, there is no satisfactory explanation of the difference in the textures of copper and a brass. The most complete study of the effects of solid solution alloying on deformation textures has been that of Brick, Martin, and Angier.¹ These investigators reported that the copper texture was retained by alloys of copper with up to 30 pct Ni, less than 5 pct Zn, less than 2 pct A1 or Co up to the limit of solubility. Copper containing more than 5 pct Zn, more than 2 pct Al, or the maximum soluble concentration of tin, silicon, or manganese showed the deformation texture of a brass. Bakarian² reported that additions of 0.20 pct Ca to 99.98 pct Mg or 2.0 pct Mn plus 0.15 pct Ca to magnesium resulted in entirely different textures than for the pure magnesium. The texture obtained for high-purity magnesium sheet showed the basal plane to be concentrated parallel to the rolling plane. The texture obtained for the alloy containing 0.20 pct Ca had the basal planes concentrated in elliptical areas whose centers were tilted 15" from the center of the pole figure toward the rolling direction. This split basal orientation was even more pronounced in the alloy containing 2.0 pct Mn and 0.15 pct Ca. Fuller and Edmunds" studied the rolled sheet textures in a zinc alloy containing 1 pct Cu and 0.01 pct Mn. The basal pole figure of this alloy contained a region of strong concentration not found in the pole figure for pure zinc.' In the above cases, results were discussed in terms of possible changes in slip and/or twinning systems and critical shear stresses. However, there were no experimental data to relate these speculations with the observed textures. It has been suggested that the deformation textures of titanium, zirconium, and beryllium could be directly ascribed to their low c/a ratios. The complex deformation mechanisms of titanium" also have been suggested to be characteristic of hexagonal close-packed metals with low c/a ratios. The purpose of this investigation was to study the cold-rolled sheet textures in binary titanium-base a alloys in which the c/a ratios were greater than, less than, and about the same as the c/a for iodide titanium. The annealed textures of these alloys will be reported in a subsequent article. Materials and Methods The chemical analyses of the metals used in this investigation are given in Table I. The alloys were prepared as 15 g buttons on a water-cooled copper block in a vacuum arc-furnace. The furnace was evacuated to pressures of the order of 10.' mm Hg, flushed with 99.98 pct pure argon, and melting was carried out in 0.2 atm pressure argon. Each button was remelted until radiographs and chemical analyses indicated complete solution of the solute elements. All buttons were forged and homogenized 72 hr at 500°C. The buttons to be rolled were ground and etched to sizes and shapes suitable for the rolling

Jan 1, 1954

By E. B. Burwell, R. H. Nesbitt

Designed by army engineers, the NX borehole camera provides the most economical method of subsurface exploration so far devised. Continuous cylindrical color photographs now reveal minor flaws in foundation bedrock not always disclosed by borehole samples. A unique projector transforms the photographs into true-to-scale images of three dimensions. ANEW milestone in the progress of photography and subsurface exploration was reached when the Chief of Army Engineers announced the development of the NX borehole camera. In obtaining continuous undistorted cylindrical color pictures of dry or water-filled borings the device surpasses all other cameras designed for photographing interior surfaces of pipes, wells, and conduits. The usefulness of the camera and the unique projector that comes with it are at once apparent. To the geological, mining, and civil engineer alike it signifies the end of guesswork with respect to underground conditions that have eluded identification by conventional core-drilling methods. To the medical profession it may well signal new progress in photo optics which could result in cylindrical color photographs supplementing the X-ray and existing photographic methods used in exploratory surgery. The pilot model of the camera has more than paid for itself as a result of its success in bringing to the surface at two dams information previously obtainable only by costly drilling of mansized boreholes for subsurface on-the-spot examination. Such bores cost as much as $200 per lineal foot as against $10 per ft for the conventional 3-in. or NX exploratory core borings from which the present invention derives its name. The NX camera can now photograph completely the interior surface of these borings. Research and development for the first camera and projection equipment were conducted by Engineering Research Associates, Division of Remington Rand, Inc., under the technical supervision of the Geology and Geophysics Branch, Office of the Chief of Army Engineers in Washington, in accordance with plans and specifications furnished by this branch. Subsequent modifications of the camera-lowering device, transporting equipment, and special field tools were made in the Corps of Engineers' Ohio River Division Laboratories in Cincinnati. The background of this development is interesting. Among the more difficult foundation problems involved in construction of concrete dams are those resulting from bedrock imperfections that escape disclosure by conventional exploration methods. Often these imperfections are of sufficient magnitude to introduce costly changes, especially when their presence is discovered only after foundation excavation is well in progress. Bedrock flaws most frequently camouflaged may be ruptures resulting from the enormous stresses to which rocks have been subjected, or they may be planes or zones of weakness related to chemical alteration and underground erosion. Obviously, detailed information on their location, dimensions, and structure is important, as these factors can affect the stability of a dam and the security of life and property downstream. Their discovery and appraisal entail extensive subsurface exploration. At present the most economical exploratory tool available to geologists and engineers is the small-diameter diamond core drill by means of which cores of the foundation bedrock are recovered for examination and testing. All too often, however, these samples fail to disclose minor foundation flaws that may be of major importance before a structure is completed. As a result, the small-diameter drillings are supplemented generally by more reliable but much more costly shafts, tunnels, or large-diameter calyx drill holes that permit the investigator to examine the foundation rocks in place. It was to reduce the number and cost of these expensive and time-consuming investigations, as well as to obtain more complete information on rock structure from the smaller borings, that the NX borehole camera was developed. That it has been eminently satisfactory in accomplishing that purpose is demonstrated by its success most recently at the Folsom dam, under construction on the American River, 23

Jan 1, 1955

By T. Federighi, S. Ceresara, C. Panseri

The results of an exploratory investigation by means of resistivity measurements at -195ºC, on the aging characteristics of a set of Al-Mg alloys with a maximum content of 7 pct Mg, are reported. Whereas for dilute alloys (Mg < 1.5 pct) only a decrease in resistivity is observed at about 100°C (due to the elimination of quenched-in vacancies), for more concentrated alloys such a decrease is preceded by an increase in resistivity in the range between -60ºand +40°C. Following some previous suggestions, this increase is considered as good evidence for clustering of solute atoms @reprecipi-tation process). From the analysis of the results it appears that, although the process is substantially similar to that observed in other Al-rich alloys, it is characterized by a strong interaction between vacancies and zones. THE process of formation of Guinier-Preston zones, i.e., preprecipitation,1&apos;2 which occurs during the aging at about room temperature in many supersaturated aluminum-rich alloys, has been extensively investigated in the past few years by several workers.2,3 owever, a very interesting question arises when the case of Al-Mg alloys is considered: although it is very easy to supersaturate Al-Mg alloys, just in the same way as for many other aluminum alloys showing preprecipitation, only very little experimental evidence was available until a short time ago in the literature supporting the existence of "natural" aging in supersaturated Al-Mg alloys (see, for example, Refs. 4 to 6). On the other hand, as is well known, the X-ray method is not useful in this case because the X-ray scattering factors of magnesium and aluminum atoms are about the same. Recently, however, Cordier and Detert7 have stressed that during the aging at about room temperature of a supersaturated Al-Mg 10 pct alloy it is possible to detect an increase in resistivity; since, as is well known, such an increase is observed also in most alloys showing preprecipitation, the above authors have interpreted their results as evidence for the formation of magnesium-rich clusters (or Guinier-Preston zones) during the aging. In the course of or ientative investigations on preprecipitation in several aluminum-rich alloys, we also could detect a similar phenomenon in more dilute Al-Mg alloys, see Instituto Sperimentale Metalli Leggeri, Annual Report, 1959, and also Ref. 8. Since the details of our observations, although of a preliminary nature, differ from those of Cordier and Detert and allow one to draw a more advanced picture of the preprecipitation process in Al-Mg alloys, we have thought it useful to report our results in this paper and to comment on them briefly.

Jan 1, 1963

By C. G. Dunn

MANY investigations have been made concerning the nature of plastic deformation and recrystallization of metals either in the form of polycrystalline materials or in the form of single crystals. However, similar work on strip materials composed of large grains as thick as the strip and having a high degree of preferred crystal orientation has not been carried out extensively. Such materials are ideal not only for studying changes in single crystals under conditions where neighboring grains exert their influences but also for obtaining the, complete change in orientation of a grain or group of grains during the course of plastic deformation and subsequent recrystallization. It is well known that first-order twins in the form of Neumann bands may be produced during the plastic deformation of ferrite crystals, particularly silicon ferrite crystals.1-3 Furthermore, it is believe that high-stress regions exist in the vicinity of these bands4a and it has been said that these banded regions would be active in recrystallization processes.5 Evidence to that effect in silicon ferrite had already been presented with photomicrographs showing recrystallization among Neumann bands during the early stages of recrystallization.5,6 Similar evidence has been published for nonferrous metals.5 Merging of fine parallel deformation bands during annealing to form twin bands of ordinary appearance in alpha brass has been offered not only as evidence of an active role being played by mechanical twins but also as evidence that such twins existed prior to recrystallization.5 Proof that such new structures are actually first-order twins, or even twins at all, must be considered inadequate except for those cases where X-ray or optical data have been obtained to confirm the necessary orientation relationship.7,8 A. B. Greninger has obtained excellent iliustrations of annealing twins in alpha irons Since his material was previously deformed by sharp hammering, a process that generally induces Neumann bands, the existence of twins following annealing might suggest that they grew from mechanical twins. In connection with transformation twins in alpha iron, Greninger reported that grains containing. three twin orientations occurred rarely and none occurred with three sets of well-developed bands.9 The work at the Pittsfield Works Laboratory on silicon ferrite in the form of large-grained strip steel deformed a small amount and annealed to partially recrystallize the grains indicates that the deformation and recrystallization processes that lead to annealing twins involve more complex processes than the formation of Neumann bands and the growth of such bands. Some groups of annealing twins are obtained which contain not only three well-developed twin structures, but also contain more than one order of twins. It is also shown, in one case, that an interrelated group of twins did not arise within one single parent grain and, furthermore, orientations of grains in the group

Jan 1, 1944

By Carls Benedicks

It is well known by every one who has had to deal with boiler tubes that these are often seriously affected by a sort of corrosion, occurring as a local pitting, that frequently causes a perforation of the tube in a rather short time. The characteristic appearance of this corrosion is shown in Fig. 1. A very similar corrosion occurs in condenser tubes (Fig. 2).l A study of the microstructure or local chemical composition of the metal (iron 70:30 brass) fails to reveal any cause for this pitting. Of course one might possibly assume the starting point to be some oxide particle, causing an electrolytic local action, but this in no way explains the continued corrosion, progressing even when the assumed oxide particle must have disappeared. By the excellent work done by the Corrosion Research Committee of the Institute of Metals, and communicated, in October, 1923, to the Institution of Engineers and Shipbuilders in Newcastle-on-tyne, stress was laid on the faot that a content of entangled air has a very obnoxious effect. In the following correspondence, Sir George Goodwin stated that "engineers have, for a long time past, regarded air as the big enemy to be fought against in all questions of corrosion&apos;72 and Kenneth Fraser pointed out that the most vital point is "the avoidance or removal of possibly the chief cause of present day corrosion, i. e. entrapped air and gases brought out of solution."3 In their reply, however, the authors state that they had "never for a moment believed that entangled air was responsible for the beginnings of corrosion, and express the opinion that the gradual liberation of air could not account for localized corro~ion."~ On consulting the many textbooks on corrosion, as that

Jan 1, 1925

By F. R. Zacher

Unmined blocks in the Sewickley seam, surrounded by worked out areas, have been found to transmit overburden oads through the interval strata to the Pittsburgh seam workings 90 ft below. Operating experiences under these conditions are described in this paper. IT has always been understood, in northern West Virginia where both the Pittsburgh and Sewickley seams are mined, that pillaring or splitting in the lower Pittsburgh seam could break the interval strata and make mining in the overlying Sewickley difficult or impossible. In some instances, however, mining methods practiced in the overlying Sewickley have reversed the problem, resulting in a shifting of the intermediate layer and destroying valuable acreages of the Pittsburgh seam below. The Pittsburgh seam in this area is, on the average, 102 in. thick. About 18 in. of low-quality head coal are left for roof protection as it is mined. Immediately over the coal are two draw slates, two rider seams, and a good hard black shale, locally called black rock, which serves as a permanent roof where draw slates and riders are taken down on certain main haulroads. In most areas mined in this seam in the past the bottom has been satisfactory but soft. However, in areas now being mined by the Christopher Coal Co. in Monongalia County the bottom is very hard and unyielding. The Sewickley seam, lying 85 ft above the Pittsburgh, will average nearly 70 in. in thickness and is overlain by massive shales and sandstones. Much of the seam is mined only down to the "sheep skin," a hard band about 4 in. above the bottom, which is composed of relatively hard shale. Fig. 1 shows a typical cross-section through the Pittsburgh and Sewickley seams in the area being discussed. The thin Redstone seam is not deep-mined but is stripped in some locations where it outcrops. Most of the Sewickley seam in the Morgantown areas has been mined out from 10 to 30 years. There are at present only two large mines and five or six others of small production operating in that seam. Mining systems in the past were generally of the room and pillar type with some full pillar extraction. This seam in the Scott&apos;s Run area was operated by as many as thirty or forty different companies at one time, during an era beset with strikes, market failures, and depressions. Most of the barrier pillars along the many main haulroads were abandoned, and blocks of Sewickley coal that would be fully recovered with today&apos;s machinery were left unmined when conditions became anything less than favorable. Consequently the mined-out Sewickley is spotted throughout with solid blocks varying in size from 100x100 ft to 200x400 ft and larger. It is these unmined Sewickley blocks surrounded by worked-out portions that have caused transmissions of roof stresses to underlying Pittsburgh mining areas, see Fig. 2, resulting in untold difficulty. As the Christopher Coal Co. operates both the Pittsburgh and Sewickley seams, accurate maps of workings in both seams are readily available and are regularly consulted. The first Pittsburgh squeeze attributed to overlying Sewickley blocks occurred in the Arkwright mine, where a shuttle car crew was developing a section consisting of seven panel headings. These entries were being driven in solid coal 3400 ft from the nearest worked-out Pittsburgh area. As mining advanced it was noticed that certain ribs began to slough and certain areas to show signs

Jan 1, 1953

By J. W. Axelson, J. N. S. Kwong, J. T. Adams, E. L. Pire, J. F. Johnson

CRUSHING has always been a major operation in the chemical and metallurgical industries, yet little is known about the theory of crushing, and today, the design of crushers is still based almost entirely on empirical knowledge and accumulated practical experience. In view of the increasing national need for the economic working of poorer grades of ores, the lack of a fundamental understanding of this unit operation hardly presents a satisfactory situation. Basic investigations of crushing have been concerned mainly with three phases of the problem, 1—the mechanism of the fracture process itself, 2— the particle size distribution of the crushed product, and 3—the relationship between the energy input and the amount of new surface produced. Probably adequate information on these phases will be required for a comprehensive understanding of the process of crushing. A bibliography covering these phases will be given as well as a review of the work done in recent years at the University of Minnesota on the relationship between the energy input for crushing and the amount of new surface produced. Fracture of Solids This phase of the crushing problem is the most fundamental because it is concerned with the actual mechanism of fracture. The main problems in a study of the fracture process are concerned with the questions of why and how fracture occurs and why there is such a discrepancy between the actual and theoretical energies needed in a fracture process. In considering the how of fracture, Poncelet40 carried Out experiments On the crushing of glass under compression. Actual photographs were taken of the glass plates in various stages of fracture-From these experiments Poncelet was able to postulate a probable mechanism of fracture. Probably because of the quantitative nature of the problem, considerable effort has been expended to find an explanation for the low tensile strength of materials and the conversely high energy for crushing. Tensile strengths are often only 1/500 of the theoretical while the energy for crushing is usually at least 500 times the theoretical. A review of the work to explain the low strength of glass is given by Weyl,52 in which the work of Griffith," Joffe,32 and Powell and Preston " are discussed. In a recent paper Seitz46 holds that the presence of flaws and fissures, as originally expounded by Griffith, is the most probable reason for the low strength of metals. Since a mass of data has been accumulated on the fracture strength of materials, several investigators have attempted to correlate this data in the form of an equation. Glathart and Preston," Taylor,40 Ponce-let,40 Machlin and Norwich," and Frederickson and Eyring,20 have all considered this problem, and the final relationship has always been that the stress for fracture is proportional to the logarithm of time. Murgatroyd30 arrived at a very similar relationship by a different method. The study of the fracture process has used many methods. Among some of these are the thermody-namic approach used by Furth21 which was based on the work of BornxL and later used by Saibel,43 and the statistical method applied by Fisher and Hollo-mon.10 In recent years metallurgists have concerned themselves with the problem of brittle .fracture as opposed to the usual plastic failure of metals. The great amount of interest shown is reflected by the number of contributors to the recent symposium on the fracture of metals.&apos; Repeated failures of welded ships during World War II have accelerated these Studies.31,50,34 The shifting of thought from an emphasis on the

Jan 1, 1952

By R. Fawley, M. A. Quader, R. A. Dodd

Stacking fault energies and annealing twin frequencies have been measured for austenitic stainless steels having a) constant 20 wt pct Cr and varying nickel contents, and b) constant 20 wt pct Ni and varying chromium contents. The effect of nitrogen has also been studied and, to a lesser extent, so has the effect of carbon. All measured stacking fault energies plot linearly against the corresponding annealing twin frequencies, suggesting that, for constant prior deformation and grain size, twinning frequency can be used to estimate stacking fault energy. Compositions low in nickel possess the lowest fault energy, while the maximum fault energy within the investigated composition range occurs mound 20 pct Cr-20 pct Ni. Nitrogen has practically no effect on stacking fault energy when increased from 0.005 to 0.05 wt pct, although Planar dislocation glide is associated with higher nitrogen contents; this may be a short-range order effect. Carbon may raise the stacking fault energy a little, and causes pronounced dislocation tangling in deformed alloys. The basically different dislocation structures promoted by nitrogen and carbon, respectively, are not associated with large variations in stacking fault energy, and the reason for the difference remains unknown. INTEREST in the dislocation configurations developed in plastically deformed alloys arises principally from a desire for an understanding of mechanisms of plastic deformation and fracture. Comparatively little basic work has been concerned with austenitic stainless steels, although considerable interest has developed in the fundamental aspects of deformation of these alloys since the suggestion that those compositions deforming essentially by planar glide are particularly prone to transgranular stress corrosion cracking.&apos; Certain facts are well-established. An 18 Cr-8 Ni steel (type 304) has a low stacking fault energy which is increased by higher nickel contents, and the structure of these steels when moderately plastically deformed changes concomitantly from planar to less planar to distinctly tangled dislocation configurations. These changes are thought to relate to changes in stacking fault energy, and considerable effort has been expended in measuring stacking fault energies of austenitic stainless steels through the determination of node curvatures2"4 even though the method has limitations due to formula approximations,5 an upper limit of measurement of 50 erg cm-2 at the most,5 and the necessity of an extremely meticulous experimental technique. Therefore, it is not surprising that only qualitative agreement between different investigators is generally to be found.&apos;-&apos; At high energies where node radius measurements are inappropriate, some use has been made of annealing twin frequencies to estimate relative stacking fault energies, although a recent critical evaluation of the method suggests that it should be used cautionsly with due regard to effects of grain size and prior deformation.6 Finally, there is a question as to whether the structures of some deformed steels are due to ordering effects. In particular, the role of nitrogen in promoting planar dislocation motion has been suggested to be due to ordering,&apos; but similar suggestions have been made for steels having low interstitial contents.4 EXPERIMENTAL Small ingots of austenitic stainless steels (-50 g) were prepared from components exceeding 99.99 pct purity, the nickel and chromium being Johnson Matthey "Specpure" metals. Carbon was added as pure graphite, and nitrogen as CrN. The components were accurately weighed, and it was assumed that the final composition was the same as the nominal composition

Jan 1, 1969

By N. K. Chen, C. H. Mathewson

Single crystals of high purity aluminum of various orientations were carefully documented after plastic extension. Special attention was given to the formation of slip lines, deformation bands, and asterism. An explanation for the geometry and mechanism of formation of deformation bands is considered. WHILE the general process of plastic extension of face-centered cubic metals is well understood,&apos; the accompanying structural changes due to slip and rotation are far more complex than ordinarily assumed. In the case of stretched single crystals of aluminum, the complicated behavior is especially marked by the formation of deformation bands, as first indicated by Collins and Mathewson." Such bands were believed by them and also by others,".&apos;7 who have discussed banded structure in general, to represent one of the important modes of lattice reorientation and to be intimately connected with the recrystallization process. Yet, despite their fundamental importance to the understanding of the mechanism df both plastic deformation and recrystallization, practically no study of the crystallographic features of these bands in their relationship to the conventional slip process has been attempted. This paper describes part of an extensive investigation on the structural behavior during plastic deformation of aluminum single crystals and is especially concerned with the characteristics of deformation bands. The investigation of the subsequent re-crystallization of deformed aluminum single crystals will be discussed in a later publication. Experimental Procedure Single crystal specimens of 99.996+ pct pure aluminum, Yz in. in diam and 5 to 6 in. in length were made both by the Bridgman method5 nd by a modified temperature gradient method.&apos;* Crystals were selected in which the lineage deviation did not exceed 2"; i.e., as determined from the back-reflection Laue X-ray photograms. These crystals were radiographed to insure that no internal porosity was present. The orientations showed good coverage of the whole area of the Taylor and Elam triangle.&apos; Orientations of the crystals were obtained by the back-reflection Laue technique described by Greninger." Crystals so selected were tapered very carefully in a lathe to obtain a linear decrease in diameter from 0.500 to 0.350 in. along a gage length of 3 in. This is considered desirable because it permits a strain gradient to be introduced in a single specimen at one initial orientation. The advantages of tapered crystals for studying the development of slip lines and recrystallized grains have been pointed out by previous investigators." The cold-worked layer resulting from the machining operation was removed by alternate etching in aqua regia and careful polishing with 2/0 emery paper until undistorted Laue back-reflection spots were obtained. A homogenization treatment at 600 °C for 24 hr was then introduced to insure the removal of residual stress. Two of the crystals so prepared were also checked by use of a glancing angle technique. Using a copper target, at 30 kv and 10 microamperes, a monochromatic beam was positioned at right angles to the axis of the crystal about which it rotated. The time of exposure was 24 hr. From these X-ray photograms, it was concluded that the present method of preparation of single crystals is satisfactory. Following this procedure, the crystals were elec-trolytically polished in a 2:l solution of methyl alcohol and concentrated nitric acid. Intermittent

Jan 1, 1952

By N. K. Chen, R. Maddin

In the extension of molybdenum single crystals at room temperature, the slip planes were found to be of the type (1101; the slip direction <111>. Theories of plasticity of body-centered cubic metals have been examined, and an explanation based upon planes of highest atomic density seems plausible to explain the plastic behavior of molybdenum single crystals. IN body-centered cubic metals, slip has been reported to occur on several planes; these are of the type {110), {112), and (123). Andrade has correlated&apos; the operative slip planes in body-centered cubic metals with the testing temperature relative to the melting point. In this correlation, if T is the absolute temperature at which slip takes place and Tm the absolute melting point, then the acting slip planes in body-centered cubic metals may be arranged as shown in Table I. However, deviations from this correlation have been reported in iron and iron-silicon alloys2 and also in molybdenum deformed at a temperature around 2400°C." In the case of iron and other body-centered cubic metals or alloys in which all three slip planes have been reported to be active at any one temperature, the crystallographic mechanism of slip is not well understood. Numerous investigators in this field have assumed that the resolved shear stress is a deciding factor. Such considerations were given by Gough in his experiments with repeated torsion of iron single crystals4 and recently by Opinsky and Smoluchowski in their study of iron-silicon alloys," and will be dealt with later in this report. A third possibility is that slip lines in body-centered cubic metals may be formed by alternate slip on nonparallel (110) planes and not by slip on any plane of higher indices. This was advanced by Elam in the case of ß brass." Since (1101 planes are planes of highest atomic density in the body-centered cubic structure, this hypothesis would be in accordance with the plastic behavior of metals in other systems, namely, face-centered cubic and hexagonal close-packed. Experimental evidence in support of this hypothesis is lacking. Nevertheless the work of Barrett, Ansel, and Meh12 has definitely demonstrated that the slip resistance on {110} planes in iron is very much smaller compared with that on (112) and (123) planes. In the present investigation, molybdenum single crystals were carefully documented as to the operative slip planes and lattice reorientations after plastic extension at room temperature. In three crystals studied, (110) planes were found to be active invariably rather than the (112) planes, which were previously reported for molybdenum at this temperature&apos; and predicted by Andrade&apos;s theory.&apos; In two other cases (112) and (123) planes seemed to operate together with (110) planes as active slip planes. However, it is believed, after careful analysis of crystallite and lattice rotations, that slip on (112) and (123) planes may be spurious. Particular reference will be made to the formation of asterism, lattice rotation, and crystallite rotation in an attempt to show that slip apparently observed on {1121} and (123) planes is actually slip on nonparallel {110} planes. Materials and Treatment Sintered molybdenum rods ? in. in diam, obtained from the Fansteel Metallurgical Co., were reported to have a purity of 99.9 pct.* One specimen, Mo-13, ¼ in. in diam, was supplied by Westinghouse Electric Corp. Single crystals were prepared, using the technique devised by Chen, Maddin, and Pond." Mo-13 was ¼ in. in diam x 7 in. long with a reduced gage section of 0.200 in. in diam x 3 in. long. The center of

Jan 1, 1952

By R. D. Nelson, F. E. Bowman

Evidence has been found for a (010)a?(111)B[102]a /,11[341]/B crystallographic relationship between the monoclinic a and B phases of plutonium during the a?B transformation. The results of an X-ray diffraction study, where [010] fiber textured a samples were used in the absence of suitable single crystals indicate that the (OIO)a planes become (111)B planes upon transformation under near-equilibrium conditions. The transformation strain along the fiber axis, which was the [010] direction, was measured dila-tomekically; the measured 8.5 pct expansion compares favorably with the 8.44 pct expansion in the in-terplanar spacing represented by the (010)a?(111)B relationship. From plots of the atom positions in each of the related planes, a correspondence was deduced between the [102]a, direction and the [341]B direction. Further analysis of the two lattice structures in the light of this relationship revealed a rotation about the [102] direction of 1 deg 43 min and a slight shearing in this same direction of about 30 min. THIS paper discusses the (B crystal structure of plutonium and the orientation relationship existing between the monoclinic a and B phases during the a?ß transformation. The kinetics of the a?ß and (3 — a transformations have received extensive consideration, and because of the crystallographic complexity of these phases, the question of the mechanisms by which these transformations occur is especially intriguing. Nelson and Shyne1 evaluated the existing experimental evidence and concluded, from the indirect evidence available, that both the ß ? a and a ? ß transformations occur by a martensitic process. Their data reveal the anisotropic nature of these transformations which suggests the existence of specific crystallographic relationships between the parent and product phases. In this paper, we are concerned with the confirmation of the existence of such a relationship resulting from the a ? ß transformation and the derivation of the crystallographic details necessary to specify it. EXPERIMENTAL RESULTS X-ray diffraction and dilatometric data were obtained from samples of a plutonium having a [010] fiber texture. These cylindrical specimens, approximately 1 cm in diam, were prepared by transforming the ß phase to a under the influence of a uniaxially applied compressive load.&apos; The [0l0] direction is parallel to the direction of applied stress; it is also perpendicular to the (010) planes. This correspondence between the plane normal and the direction having the same indices is, in the monoclinic a plutonium structure, restricted solely to this (010)[010] combination. Such a specimen was mounted in an evacuated, elevated temperature specimen holder on an X-ray dif-fractometer so that the diffracting surface was perpendicular to the fiber axis. The a diffraction pattern, Fig. 1, taken at about 115ºC, revealed the predominance of the (020) reflection. This monoclinic, P 21/rn space group yields no (0k0) reflections when k is odd. The temperature was raised slowly to 160°C to effect the complete transformation to the phase; then it was lowered to 120°C so that a diffraction pattern of the ß could be obtained near the a = ß equilibrium temperature for direct comparison with the a at the same temperature. The (3 diffraction pattern, Fig. 2, obtained under these conditions showed a (222) reflection relatively much more intense than in the random pattern. The other reflections appearing with sig-

Jan 1, 1970