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By R. E. Reed, C. J. McHargue

Single crystals of aluminurn were reduced 90 pct in diameter by extrusion at 296" and 77°K. The resultant double fiber texture had a strong (111) component with a weak (001) component. The relative anzounts of these components varied with starting orientation. Decreasing the Lemperatuve of defovtnation increased the amount of (001) texture component in crystals with inilia1 orientations near (111). The starting orientation and temperature of deformation strongly affected the morphology of the (001) component. The effects of orientation and Lemperatuve upon the distribution of the (001 ) texture component in extruded aluminum single crystals were rationalized on the basis of deformation band formation and interaction. MANY papers since the early 1920s have discussed the occurrence of the double fiber texture (111) + (001) in wires and rods of deformed fcc metals. Aside from the reporting of experimental results for the various metals under a variety of fabrication conditions, emphasis has been placed on two related problems. One problem deals with the prediction of ideal end orientations based on the glide systems and flow directions. The other problem seeks the reason why different amounts of each component may be produced by the same fabrication schedule for metals apparently having the same deformation systems. Some of the difficulties in developing a satisfactory explanation have their origin in the experiments themselves. One mistake has often been the assumption of an initially randomly oriented specimen or complete ignorance of effects of starting orientation. Recent studies by Grewen and Wassermann,1 Cullity and co-workers,2' and McHargue4 have shown that the initial orientation may be the most important factor in determining the final orientation. Hence only limited comparisons of data from different studies can be made unless all conditions of the initial material are specified. Another question has been the nature of the (001) component. Hibbards suggested that any (001) present was the result of re crystallization during fabrication. Although Fleischer6 found no (001) texture in aluminum swaged at room temperature, and Dayal 7 found that component to disappear at very high reductions, a double fiber texture has usually been reported for aluminum. Vandermeer and McHargue showed that, while some (001) in aluminum may be the result of recrystallization, much of this component has its origin in the deformation process. Ahlborn and Wassermanng found the (001) component in silver deformed at low temperature not to be the product of recrystallization.

Jan 1, 1968

By A. Raman, P. P. Singh

The constitution of the La-Rh system was studied by powder X-ray diffraction, metallopaphic, and differential thermal analysis techniques and an equilibrium diagram is presented. Eleven intermediate phases occur in the system and the crystal structural data for nine of them were determined. La3Rh crystallizes in an orthorhombic structure of undetermined type, whose unit cell is obtained by doubling the 'a; and 'c,,' edges of an FesC type unit cell. The other intermediate phases of the system are LarRh-3( undetermined structures also occur in the system. LaRh, undergoes a polymorphic phase transformation at 1240°C. LaRh3 and La2Rh7 also exhibit polymorphisnz. The phases Laah and LazRh7 melt congruently. The latter undergoes a eutectoid transformation into LaRh, and Rh at 1205°C. Laah3 is formed by a peritectoid reaction between Laah and La,Rh,,. The other Phases result from peritectic reactions between the liquid and the adjacent rhodium-rich phases. The intermediate Phases of the La-Rh system are compared with those of the La-Co and La-Ni systems. DURING the course of a detailed investigation to study the occurrence of CrB, FeB, A1B2, and related structures in the rare earth alloys it was found that much information is lacking for the rare earth noble metal systems. Although the structures of several rare earth alloys containing the noble metals at the AB and AB2 stoichiometries have been reported, the occurrence of related structures at other stoichiometries has not been studied. We have initiated a project to study the crystal structural features of selected rare earth-rhodium alloys and to map the equilibrium diagrams of representative systems with conventional methods. The results of our investigation in the La-Rh system are presented in this paper. Two phases were known in the La-Rh system. LaRh has the CrB-type structure.' LaRhz is a MgCu2-type Laves phase.z EXPERIMENTAL PROCEDURE Alloys weighing less than 1 g were prepared from commercially pure lanthanum (99.9 pct +), supplied by Lunex Company, Pleasant Valley, Iowa, and rhodium (99.92 pct +), supplied by Engelhardt Industries, Newark, N.J., in a conventional arc melting furnace under argon atmosphere. The buttons were turned upside down and remelted three times to insure homogeneity in the samples. Since negligible loss of material was encountered during melting, a chemical analysis of the alloy buttons was not undertaken. Powder specimens for X-ray diffraction studies in the as cast state were then prepared. The buttons were wrapped in thin molybdenum foils and homogenized by heating in vacuum at suitable high temperatures for more than 1 week. They were then broken into three or four pieces for annealing experiments. The pieces were wrapped in molybdenum foils and annealed at various temperatures in evacuated quartz capsules. The annealing was carried out for 2 hr at or above 1200°C, 1 day at temperatures close to llOO°C, 2 days at 1000°C, and for 1 week at temperatures below 1000°C. After annealing the alloy pieces were again broken and powder specimens for X-ray diffraction were prepared. The powders of the lanthanum rich alloys with more than 80 at. pct La were prepared by filing. The filings were sealed in molybdenum tubings and stress-relieved at 600°C in vacuum. It was not deemed necessary to stress-relieve the powders of the other alloys, since the alloys were very brittle and were ground easily. POWDER X-RAY DIFFRACTION X-ray diffraction photographs of powders (-325 mesh size) of the alloys in the as cast and annealed states were prepared in a Guinier-de Wolff focussing camera with copper K, X radiations. These patterns were studied to identify the stoichiometries and the crystal structures of the intermediate phases. The lattice parameters of the phases were calculated after minimizing the differences between the observed sin2 6 values, calculated from the diffraction angles 8, and the sin2 8 values, calculated using approximate lattice constants obtained from a few lines. These differences were minimized manually to less than 0.0005. The latLice constants are judged to be accurate to *0.005A for values less thp about 10A and to k0.01~ for values greater than 10A. The relative intensities of the lines were calculated using a computer program written by Jeitschko and Parthk.~ No attempt was made to refine the atomic positional parameters in the phases. METALLOGRAPHY The phase equilibria in the investigated alloys in the as cast and annealed states were also studied by metallographic examination. The polished specimen surfaces were etched with 10 pct picric acid in alcohol (alloys up to 25 pct Rh), concentrated picric acid (from 25 to 37.5 pct Rh), 2 pct nital (40 to 50 pct Rh), 10 pct nital (from 50 to 66.7 pct Rh) and with concentrated 48 pct HF for the other rhodium-rich alloys. Selected microstruture~ were then photographed using a Po-laroid Land camera. THERMAL ANALYSIS Differential thermal analysis of the alloys was carried out in DTA-668 Stone differential thermal ana-

Jan 1, 1970

By R. F. Mehl, J. K. Stanley

THE recrystallization of cold-worked metals is studied ordinarily by determining the temperatures required for complete recrystallization to occur within a given arbitrary time period, usually within 15 min. to 2 hr. The results obtained are conventionally assembled in a three-dimensional diagram relating this temperature to the amount of cold-work (percentage deformation) and the resultant grain size, though sometimes other modes of representation are employed.1 Such three-dimensional diagrams are not complete, for the number of variables exceeds the number represented. The process of recrystallization is characterized by a time-rate, a fact that is not portrayed upon ordinary recrystallization diagrams. It is known that recrystallization consists in the formation of new nuclei and the growth of these nuclei, a process that continues until the matrix is completely consumed and the original set of grains entirely replaced by the new. This knowledge is old.2 Truly heterogeneous reactions invariably proceed by nucleation and growth; the decomposition of eutectoid solid solutions provides an example about which much is known,3 and freezing, originally studied by Tammann, provides the classic example. Although recrystallization is not a heterogeneous reaction .within the orthodox meaning of the term, the mechanism of the process is the same. As in eutectoid inversion and freezing, it should be possible to study the rate of isothermal recrystallization as a function of time, and it should also be possible to analyze the isothermal reaction curve into the component rates of nucleation and growth that determine them.* Other methods, not employing the rates of nucleation and growth, have been used to represent the rates of such reactions, but these in general have been artificial and have not led to progress.4,5? The theory of recrystallization is not in advanced state; if recrystallization data were obtained, in terms of the rate of isothermal recrystallization, and in terms of the rate of nucleation and the rate of growth of recrystallized grains, and in particular if the temperature coefficients of these rates were measured accurately, an adequate theory should be more readily developed.27 Such an effort requires proper experimental technique, but especially requires suitable methods of mathematical analysis to convert measurements made upon a surface of polish into values for the rate of nucleation and the rate of growth and to combine these into a rate of isothermal recrystallization. These methods are now

Jan 1, 1942

By D. J. McAdam

In a recent survey of the literature, the author has found evidence incompatible with prevalent views regarding the technical cohesive strength and yield strength of metals. Some of the evidence regarding the technical cohesive strength was presented in a previous paper,21 and views were presented that are believed to be in better accord with the evidence. In that paper, it was shown that the technical cohesion limit (contrary to prevalent views) varies with the combination of principal stresses,t and that the influence of plastic deformation on the technical cohesive strength is not in accordance with prevalent views. In this paper, consideration will be given to both the technical cohesion limit and the yield stress as affected by the combination of principal stresses. Evidence will be presented that the variations of the yield stress and the technical cohesion limit with the combination of principal stresses are not in accordance with prevalent views. Consideration will also be given to ductility and work-hardening capacity, which depend on the initial difference between the technical cohesion limit and the yield stress. By the "technical cohesion limit" is meant the mean stress at actual or poten- tial* fracture. An example of a rechnlcal cohesion limit is the "true" breaking stress of a tension-test specimen, the breaking load divided by the sectional area at fracture. The cohesion limit thus obtained for a ductile metal, however, has been greatly altered by the plastic deformation. It is important to consider the cohesion limit of the metal prior to this deformation, and the variation of the cohesive strength with plastic deformation. The relation between the cohesion limit and the corresponding flow stresst determines the ductility of a metal and the energy that it will absorb before fracture. In this paper, as in the previous paper,21 tensile stresses will be viewed as positive and compressive stresses as negative. The algebraically greatest principal stress will be designated S1, the least principal stress will be designated S3, and the intermediate principal stress will be designated Sz In the previous paper, attention was confined to stress combinations with two of the principal stresses equal (with the intermediate principal stress S2 equal eithel to S3 or 31). In this paper, consideration will also be given to the effect of stress combinations with all three principal stresses unequal. In the earlier sections of the paper (pp. 3I I to 324), attention will be confined to stress combinations with

Jan 1, 1942

By Granville S. Borden, Frank H. Madison

The fruits of industry are divided between capital, labor, and governments. Capital takes its redemption and remuneration through profits or dividends; labor takes its share through wages; governments take their share through taxes. Each must have a just share, and the question of what is a just share and the question of who shall make the allocation are vital current political and economic issues. Certain political philosophies con- tend that government should take all the fruits and that some dictator or politician should have irrevocable discretion and power to distribute the fruits. Believers in democracy and free enterprise adhere to the view that capital is entitled to redemption and fair remuneration; that labor shall have a fair wage; and that governments shall make equitable levies. But since all human beings, both individually and in groups, are inherently selfish, there are many variable views in regard to the question of what ratios shall be applied in a division of the fruits and who shall have the power to fix these ratios. Perhaps all parties should devote more energy to means of increasing the crop, and not so much time to quarreling over a division of the existing crop. Moreover, all three parties must guard against killing the goose that lays the mineral eggs. For example, if the tax burden becomes oppressive and the taxpayer refuses to carry it, then business stagnates, with consequent unemployment, reduction of profits, and curtailment of government revenue. If a producer assumes the increased tax burden, marginal operations must be closed down, and, again the consequences are increased unemployment and diminished government revenue. Quite often an increase in a rate of tax results in a reduction of government revenue. Perhaps the worst detrimental feature of selfishness and greed on the part of any one of the three distributees of the fruits of industry is the retardation of expansion of the enterprise. Growth means new jobs, more government revenue, better trade, and better standards of living. Explorations and discoveries have resulted in tremendous contributions

Jan 1, 1964

By F. R. Hensel

PURE silver and silver-lead alloys have been studied as to their suitability for bearings.1-8 A review of the properties of thallium and the silver-thallium constitutional diagram was made by the author to analyze the possibilities of silver-thallium compositions for antifriction materials.† The silver-rich end of the diagram as reported in the literature9-11 was found to be sketchy and it was necessary to carry out considerable experimental work to arrive at definite conclusions. Some of the results of this work are reported in this paper. TEST MATERIALS In the beginning of the work, only fused alloys were investigated. Later, research [TABLE I. Composition of Fused Silver thallium Alloys Tested for Antifriction Properties PERCENTAGE ALLOY No.OF THALLIUM 15330.48 153406 231 HT2.04 2312.10 232 HT3.83Balance 2324.07Silver 233 HT5.86 2336.38 2349.84 234 HT10.07] was carried out on electroplating methods and diffusion processes. In preparing the fused alloys, care was taken to exhaust the toxic fumes caused by the thallium content. The compositions of the series of fused alloys are listed in Table I. The silver-thallium alloys were melted in clay-graphite crucibles and cast into preheated steel molds of 3/4-in. diameter. In the cast condition, they showed a cored structure, as indicated by the micrographs of Figs. I to 3. The etching reagent used was a mixture of 2 grams K2Cr2O7, 8 c.c. H2SO4, and 100 c.c. H2O. The cored structure is unsatisfactory for the type of corrosion resistance required for bearing applications, and homogenizing experiments were carried out at various temperatures, the results of which are shown in Figs. 4 through 8. The alloys with a lower thallium concentration, which were heated for 2 hr. at 525°C. show an almost completely homogenized solid solution type of structure. With higher thallium concentration, the homogenizing temperature was dropped to 475°C. to eliminate the formation of a liquid phase. It is evident that heating for 2 hr. at this temperature did not result in complete elimination of the cored structure. As would be expected, cold-working of the cast structure is an expedient in establishing equilibrium conditions. The fine homogeneous grain structure shown in Fig. 7 corresponds to the cold-worked area under the Rockwell ball penetration when the hardness was taken before heat-treatment. A detailed structural study was made on a cast silver-thallium alloy containing 3.67 per cent thallium. Three micrographs (Figs. 9, 10 and 11) show the transition

Jan 1, 1945

By Margaret G. Scroger, Luke L. Y. Chang, Bert Phillips

Using the sealed-system technique, isothermal sections at 1000°, 1400°, and 1700°C for the system Co-W-O hare been determined. The equilibrium data were obtained by microscopic and X-ray diffraction examination of samples winch had been equilibrated in pellet form fit sealed capsules and subsequcelly quenched. The only ternary phase in the system is CoWO, which mells congruentlly at approximately 1280°C. At 1000°C, this ternary phase exists in equilihrinm with all the allows and oxide compounds in the hounding binary systems. At 1400°C, liquid forms in thin oxygen-rick region near the CoWO4 composition and only one phase assemblage in the system (W + WO2 + H'5CV7) contains no In order for tungsten to be used to its full potential as a refractory metal. it is necessary that coatings be developed which will preclude oxidation of the metal. It is true that coatings now in use do protect tungsten for short periods at temperatures near 1700 °C. but no coating has been developed which protects adequately at or above this temperature where the use of tungsten is most required. This situation exists despite the great effort in recent years to develop such coatings. and it has become evident that success in this area can be achieved only after a more fundamental knowledge of the thermochemical and therrnophysical propertics of tungsten oxidation and tungsten -metal oxidation is obtained. In our 1aboratories. a project liclrriel. At 1700°C, tungsten is the ouly solid phase which is stctble, and a large range of compositions melt to two liquids, one metallic and one oxidie. The sealed-system method of etjuilibra lion prevents the determination of the partial pressure of oxygen and the total gas pressure orer the condensed phases at equilibrium. This data is not available and can be obtained only by some other method. We can generalize on these unknowns and indicate that both partial pressure of oxygen and total pressure of the gas phase must increase as compositions change from the W-Co binary toward the oxygen apex and as temperatures increase. has been initiated which is directed at providing some of this fundamental information through a study of metal-tungsten-oxygen reactions and equilibria and the properties of any intermediate phases which result. This paper is based on a part of this program in which the system Co-W-O was studied between 1000° and 1700°C. I) RELEVANT INFORMATION FROM THE LITERATURE A) The Notable Lack of High-Temperature Work on Metal-Tungsten-Oxygen Systems. It should be pointed out that there has been no published report (of which we are aware) which deals with the study of a metal-metal-oxygen system at 1700°C. Also, there has been no report on a complete metal-luugsten-oxgen system at any temperature. As a matter of fact. before this program started. The

Jan 1, 1965

By Michael P. Robinson

It has been suggested that streaming potentials are not nomlally logged because the streaming potentials known to be generated across mud filter cakes are substantially cancelled by streaming potentials generated in contiguous shale beds. Laboratory measuments of streaming poterztials acrons hand specirnens of shale have been adduced as evidence to support the suggestion. The experiments show that streanzing potential depends on the resistivity of the equilibrating solution and on pressure drop. Meizsurenlents were made perpendicular to the bedding planes of the shales. The mechanism of streaming potential generation in shale which the suggested cancellation theory iniplies has been critically examined. Conclusions reached are tested by comparing self potentials logged before and immediately after changing the resistil~ity (but not the weight) of mud in a borehole. The resultant SP change is found to equal the theoretical electrochetnical potential change, i.e., cancellation of .strearning potential is independent of mud resistivity. Since streaming po/entials generateal across mud filter cakes depend critically upon mud reistivity, shale .strearlling potential in situ must he equally sensitive. Using plausible values for the pemleahility, porosiy arzd compressibility of shale, the van Ererdingerl-Hurst approach shows that the pressure drop between horehole pressure and formation pressure in a radially honzogeneous shale is unlikely to coincide with the zone of filtrate invasion. If, perchance. this coincidence occurs with one mud a chatzge 10 a second must result in non-coincidence, i.e., a .YP logged in the second mud should show a streaming potenlial. The dilernrna may he overcome by postrrlating: (I) perrneabilities of shale7 parallel to their beddine planes greatly exceed their cross-bedding permeabilities; (2) a resistivity-sensitive mud cake of low permeability forms either on the face of, or probnhly in the lanlinations of, many shales. These postulates are examined with reference to the pres.sure and SP histories of the McClosky and Aux Vases zones in the Illinois Basin and some srcpport for them found. It is concluded that a definitive solultion to the streaming potential problem does not exist. INTRODUCTION Recently Gondouin and Scala' published a contribution to the streaming potential problem which is notable particularly for the elegance of its experimental techniques. The conclusions drawn by Gondouin and Scala from their studies amplified those previously suggested by Schehck'. In essence they suggested that the streaming potentials known to develop across mud filter cakes" are not normally observed on SP logs because they are largely cancelled by streaming potentials generated in adjacent shales. This suggestion appears so reasonable that many workers in the logging field will be prone to accept it. We incline strongly to the view that the suggestion is basically correct. Nevertheless, it is important to realize that the cancellation hypothesis in the form put forward by Gondouin and Scala does not appear to meet all the known experimental facts. It is the principal intention of this note to call attention to what appears to be a flaw in the relevancy of the experimental evidence adduced by Gondouin and Scala and to suggest an alternative hypothesis to remove it. The consequences of the new hypothesis are compared with field data which, we believe, are unique. A second but not less important purpose of this note is to emphasize that even the new hypothesis does not fully explain all field observations. Hence we conclude that further study of the streaming potential problem is required before it can be said that a definitive solution to it has been attained. The bare experimental bones of laboratory data on streaming potentials appear to be as follows. 1. A streaming potential is developed across a mud filter cake. This streaming potential is virtually independent of cake thickness but depends on the mud resistivity. The relationship is of the form, 2. A streaming potential is developed across a shale. This potential is independent of shale thickness but depends on the resistivity of the permeating solution if the shale has been first equilibrated in this solution. The relationship is of the form, The relationship in (1) was determined first by Wyllie3, confirmed by Gondouin and Scala and qualita-

By H. H. Waller, D. L. Monroe, D. P. R. Smyth

ORIGINALLY slusher drift development was conventional, advancing the drift full 10x13-ft size at 6 ft per round. This proved dangerous and costly because the weak fractured rock of the orebody cannot be left raw for any length of time without its reacting with air to cause spalling. With this method of drift development ground conditions were worsened by many blasts and the length of time it was necessary for the drift to remain open. The problem was practically solved by driving 5x7-ft pilot drifts, then enlarging to 10x13-ft raw size by slashing with 50 to 60-ft diamond drill holes. Calculated explosive charges are used to avoid wall and roof shattering. Still more recently jacklegs and sectional steel replaced the diamond drills, resulting in the same work for less than half the cost, see Figs. 1 and 2. Crosscuts at right angles to the main haulage drifts, which cross through the center of the ore-body, divide the ore into 200-ft blocks. Eight slusher drifts on 50-ft centers draw from one block. The slusher floor is 11 ft above the rail, inclined at 1 pct grade, and the top is 18 ft below the undercut. The finished slusher drift is 102.5 or 120 ft long, depending on whether it is north or south of the crosscut. Drawpoints are staggered on each side at 17.5-ft centers. The completed drift is concreted and has an inclined ramp and a vent at the back. When crosscuts and gangways have been excavated and concreted and slusher stations installed, slusher drift development is carried out in 10 clear-cut stages: 1—Pilot drift is driven full length, 5x7 ft. 2—Station is excavated at half-way point in pilot drift. 3—Pilot drift is enlarged to slusher drift size by drilling. 4—Lower part of slash is blasted and mucked for floor concrete. 5—Floor is poured. 6—Inside upper half of slusher drift is blasted, mucked, and concreted. 7—Outside upper half of slusher drift is blasted, mucked, and concreted. 8—Drawpoints are excavated. 9—Drawpoints are concreted. 10—Concrete ramp and sheave protection recess are constructed at back end of slusher drift. A 5x7-ft pilot drift is driven full length on the center line of the proposed slusher drift using a jackleg and detachable bits on 7/8-in. hexagon steel. A 10-hp air-driven slusher and 36-in. scraper pull muck directly into cars in the main haulage crosscut, see Fig. 2. Operations are on a three-shift basis. One man mucks the previous round, drills, and blasts his round in a shift. Average advance is 5 ft per shift. When the pilot drift is completed, 7 ft of drift in the center section is enlarged to full raw slusher drift size of 10x13 ft, for the set-up to drill long slash holes. The pilot drift is slashed to full raw size from this station with jacklegs and tungsten carbide chisel bits on 7/8-in. sectional steel. Crews line the holes without trouble, working from controls established by the engineers, and the pattern of holes for the slash is arranged to give the least shattering effect

Jan 1, 1955

By Oscar E. Harder, H. A. Grove

It is now just a quarter of a century since Fred W. Taylor§(23) pub-lished his classical paper On the Art of Cutting Metals, describing his researches in which he, in cooperation with Maunsel White, had developed what is essentially the modern method of heat-treatring high-speed steels. Regarding this research, Dr. Sauveur,(22) in his address on Steel Wizards, Past and Present, at the annual banquet of the American Society for Steel Treating in Philadelphia in 1920 said, "The discovery of high-speed steel, or if you prefer, of the treatment imparting high-speed properties to certain steels by Taylor and Maunsel White I am inclined to consider as our one epoch-making contribution to the metallurgy of steel." It was at about that same time (1900) that Taylor exhibited high-speed steels, heat-treated by the improved process, at the World's Exposition in Paris and attracted much attention by showing that the steels would operate at such feeds and speeds that the tips of the tools were visibly red hot. This observation must have started metallurgists to wondering regarding the relative hardness of the various steels at red-heat temperatures. Considerable work has indeed been done in an effort to measure the hardness of high-speed steels and related materials in the red-heat range, using various methods. The result's, however, have been only partly satisfactory, owing to experimental difficulties. The present investigation had as it's immediate inception an address by J. V. Emmons, of the Cleveland Twist Drill Co., to the Columbus Chapter of the American Society for Steel Treating in 1931, in which he deplored the fact that comparatively little was known with regard to the hardness of high-speed steels at elevated temperatures, in spite

Jan 1, 1933

By A. G. Preban, R. J. De Angelis, J. B. Cohen, P. Gordon, D. C. Brillhart

Changes in stored energy, resistivity, density, X-ray line broadening, and dislocation arrangements (from transmission electron microscopy) have been mensured on copper specimens shock-loaded at 75 to 435 kbars. The stored energy in shock-loaded copper and changes in density and resistivity are larger than those after conventional deformation. A high density of twins makes an appreciable contribution to all the measured quantities. Dislocation densities rise to values of 1011 lines per sq cm at the highest pressure. The resistivity, density, and X-ray In a study of the effects of shock loading on copper, changes in stored energy, resistivity, density, hardness, lattice parameter, line broadening, and dislocation structure have been examined after subsequent annealing of specimens shocked to various pressures. This paper presents certain aspects of the data, taken by two groups, one at the Illinois Institute of Technology, and the other at Northwestern University, and compares the changes in these properties, and what they indicate about the effects of shock deformation. measurements yield dislocation densities in substantial agreement at all pressures, after correcting for the effects due to the twins. At the higher pressures, electron microscopy yields dislocation densities that are lower by a factor of about 5 to 10. Dislocation densities from stored -energy measurements are high; this finding can be rationalized by the consideration that the deformation twin boundaries present are largely incoherent. There is no evidence for homogeneous dislocation nucleation during shock loading. Shock loading any metal or alloy, such as copper, causes a large increase in yield strength, with dimensional changes of the order of only a few percent."' Other properties have not been extensively studied, but this result clearly indicates the importance of examining the substructure. Previous studies of the substructure of conventionally deformed specimens have been made using electron microscopy,3-6 X-ray diffraction,7,8 and etch pitting.9,10 The dislocation density can reach values of about 10" per sq cm; and the dislocations are arranged in cells approximately 0.15 in size. The substructure varies widely, however, from point to point in the specimen. The substructure of explosively loaded copper consists of cells of about the same dimensions6,11 but their size is more uniform.6 Profuse twinning occurs12,13 and the twins are often too small to be detected optically.6,13 The boundaries outlining these twins appear to be incoherent.14 There is now ample evidence 13-15 that twinning is a strengthening mechanism in fcc metals but this is only part of the strengthening effect in shock-loaded copper.* From the twinning systems which operate in shock-loaded single crystals,13 and from changes in texture on shock loading,17 it has been shown that both the

Jan 1, 1968

By Hans Lundberg

FOR the last 20 years, geophysical methods have been in practical use. The results accumulated during this time have been broad enough to enable us to formulate some general rules for conducting geophysical surveys to the best advantage. It is a well established fact that electrical and magnetic methods have aided the geologist to outline bodies of ore, to differentiate between various types of rocks hidden from view and also to determine the thick-ness of the layer of drift that covers the bedrock. Fundamental require-ments for a successful solution of a geophysical problem is the existence of physical differences between the various formations under investiga-tion. If such conditions exist, our problem becomes a physical, mathe-matical and geological analysis of the effects of these differences upon the structure of an arbitrarily produced, or already existing, field of force. Experience has shown that reliable and consistent results in analyzing such fields of force are obtained only if: (1) the methods and apparatus used are scientifically sound and designed to provide information with a maximum amount of accuracy, reliability, sensitivity, simplicity and speed of operation, and (2) the interpretation of the results of such an investigation is guided not only by rigid mathematical rules but also by common sense, based upon sound geological knowledge and imagination. Today, a number of methods are based upon sound principles, are highly sensitive and are well suited for field work. The instrumental developments have resulted in self-contained instruments, which permit the determination of field ratios and intensities with great accuracy. Magnetic surveys have almost become standardized. The electrical methods, mostly based on the determination of potential-drop ratios, or electromagnetic field ratios, are built as bridge circuits and it is easy to determine amplitude, as well as phase relationship of the various electrical and electromagnetic fields.

Jan 1, 1938

By D. E. Thomas, C. E. Birchenall

ALTHOUGH Eoltzmann gave a mathematical solution for the diffusion equation (for planar diffusion in infinite 01. semi-infinite systems only) in 1894 allowing for variation of the diffusion coefficient with a change in concentration, it was not until 1933 that this solution was applied to an experimentally investigated metallic system. The calculation was carried out by Matano' on the data obtained by Grube and Jedele3 for the Cu-Ni system. Since that time concentration dependence of the diffusion coefficient has been demonstrated for many pairs of metals. However, the nature of this dependence has never been fully elucidated. Many investigators have suspected that these variations could be related to the thermodynamic properties of the solutions, one of the earliest explicit statements being contained in a discussion of irreversible transport processes by Onsager' in 1931. Development along these lines has been greatly retarded by the lack of reliable data on the variation of tliffusivity with concentration, the paucity of the thermodynamic data for the same systems at the same temperatures and compositions, and an incomplete understanding of the relation of the thermodynamic properties of the activated state for diffusion to the bulk thermodynamic properties. The last factor has been discussed by Fisher, Hollomon, and Turnbull.5 In many instances where data exist, it is difficult to know which are acceptable. This problem probably applies more strongly to diffusion data than to activity measurements. For instance, four sets of observers"-" have reported self-diffusion coefficients for copper. The average spread between extreme results is a factor of about four, though the individual sets of data are self-consistent to about 20 pct. Thus one or more factors are out of control, at least in these experiments, making estimates of internal error unreliable. The most reliable diffusion data in most systems have resulted from the use of welded couples with a plane interface from which layers for analysis are machined parallel to the interface after diffusion. The layers are analyzed, and the result is a graphical relation between distance and concentration, usually called the penetration curve. Given the same set of analytical data and distances and following the same procedure in computation, different observers will generally produce diffusion coefficients which vary appreciably, especially at the extremes of the concentration range. Experiments must be carefully designed so that the precision is good enough to answer a particular question unequivocally. In the first calculation of the dependence of the diffusion coefficient on concentration in the metallic solid solution Cu-Ni, Matano found that the coefficient was insensitive to concentration from 0 to 70 pct Cu, after which it rose more and more steeply to some undetermined value as pure copper was approached.' The same behavior was reported for Au-Ni, Au-Pd, and Au-Pt.* The data used were those of Grube and Jedele which were very good at the time, but are not considered particularly good by present standards. Furthermore, the method of calculation makes the ends of the diffusion coefficient-concentration curve unreliable. For better reliability, the high copper end of the curve has been checked by incremental couples, where the concentration spread is 67.7 to 100 atomic pct Cu. The implication of the curves calculated by Matano was that diffusion is very concentration sensitive in one dilute range of this completely isomorphous system and hardly at all in the other. Matano's result is confirmed. Later Wells and Mehll0 published data on diffusion in Fe-Ni at 1300°C, which represent a thorough test of the shape of the concentration dependence curve. They ran couples with the following ranges of nickel concentration: 0-25 pct, 1.9-20.1 pct, 0-20.1 pct, 20.1-41.8 pct, 0-99.4 pct, and 79.3-99.4 pct. Although the trend of the data indicates an S-shaped concentration dependence, their curve was drawn to the pattern set by Matano. Their original data have been recalculated for the 0-99.4 and 79.3-99.4 pct couples. Wells and Mehl's points and two independent recalculations from the raw data are plotted in Fig. 1. What appears to be the best curve is drawn through them. This curve shows little sensitivity to composition in both dilute ranges with a strong dependence at intermediate composi-tions.? Similar experiments on the Cu-Pd system are reported here at temperatures where solubility is unlimited. These lead to the same type of concentration dependence for the diffusion coefficients as was found upon recalculation of the data for the Fe-Ni system. Experimental Procedure Cu-Pd: The concentration dependence of the diffusion coefficient may be determined by the use of

Jan 1, 1953

By C. E. Lundin, D. T. Klodt

The binary alloy systems, Y-Ti, Y-Zr, and Y-Hf, have been investigated throughout their entire composition regions. There is no compound formation in any of the systems, and each system is characterized by a single eutectic reaction. The eutectic compositions and temperatures are as follows: A eutectoid reaction pct Y and 870°C occurs in the Y-Ti system, whereas a peritectoid reaction,: pct Y and 880°C occurs in the Y-Zr system. Peri-tectic-type reactions at temperatures above the eutectic levels are postulated for the yttrium and hafnium transfovmations. The development of the technology of yttrium has been given considerable attention during the past few years, and studies of binary phase equilibria have, of course, taken a prominent position in this development. In many respects yttrium, in the third group of metals of the periodic table, is similar to the adjacent group of metals, titanium, zirconium, and hafnium, and the knowledge of the phase relationships of yttrium with these metals is basic to their technology. MATERIALS AND EXPERIMENTAL PROCEDURES Materials. The metals for this investigation were supplied by the General Electric Co., Aircraft Nuclear Propulsion Department. The yttrium was in the form of an arc-melted ingot, and the other metals were in the form of high-purity, iodide-Process crystal bar. Table I lists the purities of these materials. Alloy Preparation. Melting was done by conventional techniques in a nonconsumable electrode arc furnace in an atmosphere of purified argon. Melting conditions for each binary system were the same. Each alloy button was inverted and remelted several times to assure homogeneity. Accurate weights of the charges and resultant alloy buttons were obtained to indicate deviations from intended compositions. No chemical analyses were obtained since melting weight losses were consistently in the range of 0.1 to 0.2 pct of the total weight. 10- or 20-g buttons for each 5.0 wt pct composition increment were melted to survey the three individual alloy systems. Additional alloys differing in composition by 1.0 or 0.1 wt pct increments were also melted to study selected regions of the systems. Metallograpllic Techniques. Standard metallo-graphic techniques were followed for mounting and rough grinding. Preliminary polishing was accomplished using 6-u diamond paste as an abrasive on a Metcloth Lap. Final polishing was done on a Microcloth-covered wheel using 1-u diamond abrasive paste. Purified kerosene was used as a lubricant for both polishing stages. • sothermal- Annealing. Alloys were sectioned for as-cast structlure examinations and then homogenized in preparation for isothermal-annealing treatments. Homo{:enization was accomplished by cold pressing the alloy buttons followed by 72-hr anneals at 1100c. The alloys were encapsulated in Vycor or quartz for the homogenization treatments or for isothermal anneals. Resistance-wound or resistance-element tube furnaces were used for the annealing treatments. The homogenized alloy buttons were cold rolled until cracking occurred or until a -in. specimen thickness was obtained. Small -in. square) specimens for the isothermal anneals were then sawed from the alloys. Each specimen was wrapped in tantalum foil before being sealed in the capsule. Temperatures during the anneals were controlled The time at temperature necessary to equilibrate the structures during the anneals was determined for each alloy system by holding triplicate specimens of alloys at a constant temperature for three different periotls. The specimens were quenched and examined microscopically to determine the number and amounts of phases present in the micro-structure as a function of time. Melting Studies. Eutectic temperatures of the three alloy systems were established from the results of incipient-melting studies conducted on as-cast alloys. Specimens to be melted were suspended on a tungsten wire inside a graphite cylinder placed in a glass vacuum chamber. An optical pyrometer was used to follow the temperature of the specimen as it was inductively heated in a high vacuum. The temperatures were corrected for emissivity losses by standardizing the pyrometer with known-melting-point metals. Accuracy of the temperature measurements is estimated to be + 10°C. The melting point of the yttrium was determined to be 1550°C by this technique. The invariant-temperature levels were also checked by an anneal-quench technique. This technique consists of annealing a series of

Jan 1, 1962

By Albert E. Reed

The development of the modern highspeed vibrating screen, together with the increasing availability of long-lasting stainless-steel screen cloth for relatively fine-mesh separations, means that more screen circuits will be used in the future. It seems pertinent therefore to examine a few characteristics of screen-circuit products, and especially the points of difference as compared with products of classifier circuits. The most noticeable change that takes place when substituting screening for classification in a grinding circuit is a substantial drop in the circulating load, owing to the greater ease with which the "fines" leave the circuit. Where circulating loads of 500 to 600 per cent are considered fairly normal in classifier grinding circuits, the grinding circuit equipped with screens for equivalent work will have a circulating load of only 20 to 75 per cent. As a result, the products of screen circuits will show a minimum of overgrinding or sliming as compared with a classifier circuit. The desirability of this characteristic will naturally vary according to the metallurgical requirement in each case. Because screen separations make a sharper division according to particle size, the products of screen circuits can be expected to show greater crowding of particles near the point of separation. These changes in product characteristics are so important that other factors, both in the grinding and concentration processes must be adjusted if the operator is to take full advantage of the increased tonnages and higher efficiencies available to him through the use of the screen circuit. One factor that too often is overlooked when changing to screen circuits is the importance of adjusting the ball-mill or rod-mill charge, to compensate for: (I) increased tonnage of new feed, (2) decreased circulating load, or (3) more granular character of feed. These same factors should also be taken into consideration for their effect on the various types of concentration or flotation equipment that may follow the screen circuit. Over a period of years, this equipment has generally been designed and adjusted to the product of the classifier, therefore it seems obvious that some research on the adjustment and adaptation of concentrating and flotation equipment to screen-circuit products will be most profitable to the operators. To illustrate the points mentioned above, the results from three recent installations follow, with data on the comparative characteristics of screen and classifier circuits. A range of separation from 10 to 48 mesh has been chosen, using two general types of screening equipment and two types of classifiers. The first of the installations of which the results are to be examined is that of an iron-ore installation where eastern magnetite was being concentrated prior to sintering. Fig. I shows the original grinding circuit, using classification, while Fig. 2 shows the same circuit equipped with a

Jan 1, 1947

By Terkel Rosenqvist

THE desulphurization of molten iron and steel is a very complicated process. One way to arrive at a better understanding of this process is to break it down into several simpler chemical processes that can be studied individually in the laboratory. For a study of the different factors that influence the equilibrium distribution of sulphur between liquid metals and slags, several simpler equilibria may be investigated. One very important subject is the determination of the escaping tendency of sulphur in the liquid metal and its dependency on temperature and composition of the melt. Several papers in this field have recently been published.', ' Another subject is the study of the sulphur capacity of the slag. A molten slag is indeed complex, and even if sulphur distribution data for a large variety of molten slags may give empirical data about their desulphurizing power, the importance of the individual components is still not quite clear. It is accepted generally that lime is the most important desulphurizing component in the slag. The present investigation has as its purpose to study the desulphurizing power of lime in its standard state, and to provide a basis for thermodynamic calculations of the desulphurizing power of various lime-containing slags. The standard state of lime at steelmaking temperatures is solid calcium oxide, CaO. It can react with sulphur to form solid calcium sulphide, CaS. The relative stability of calcium oxide and calcium sulphide is expressed by the free energy of the reaction: 2Ca0 (s) + S1 (g) = 2CaS (s) + O2 (g) The existing free energy data for this reaction, listed by Kelley5 nd Osborn,' are uncertain to about 10 kcal and are of limited value for a calculation of equilibrium constants. Under the conditions prevailing in a melting furnace, the sulphur pressure may be expressed conveniently by the ratio H,S/H2 and the oxygen pressure by the ratio H,O/H, (or CO,/CO). The desulphurizing power of calcium oxide may, therefore, be studied by the reaction CaO + HIS = CaS + H2O. A study of this reaction may be complicated by certain side reactions: Water vapor and hydrogen sulphide may react. to form sulphur dioxide, and calcium sulphide may be oxidized to calcium sulphate. A thermodynamic calculation shows that these side reactions will be suppressed to insignificance if the equilibrium is studied in the presence of an excess of hydrogen. The apparatus used is shown in Fig. 1. About 10 g calcium oxide and 20 g calcium sulphide (laboratory qualities) were intimately mixed, and some water was added to make a thick paste. The paste was put into a thimble of zirconium silicate, which was placed within the constant temperature zone of a furnace, and capillary refractory tubes were attached in both ends. After the mixture had been heated in dry hydrogen at 1000°C for several hours all Ca(OH), and CaCO, had decomposed and CaSO, was reduced, so only CaO and CaS remained in the thimble forming a porous plug. The mixture was examined by X-ray diffraction after the initial reduction in dry hydrogen as well as after the subsequent experimental runs up to 1425 °C. It was shown that crystalline calcium oxide and calcium sulphide were always present together in about equal amounts. The unit cell edges were found to be 4.80A for CaO and 5.68A for CaS in good agreement with existing literature values." This shows that the mutual solid solubility is very small, and that the compounds are present in their standard states. Purified hydrogen was passed through water sat-urators kept at constant temperature in a thermostat bath. The amount of water vapor saturation was checked by means of a dew point method, not shown on Fig. 1. The gas mixture was passed through the capillary inlet into the furnace, where it was sifted through the porous plug of calcium oxide and calcium sulphide. The hydrogen sulphide present in the outgoing gas was absorbed in a zinc acetate solution and the hydrogen was collected over water. When one liter of hydrogen had been collected, the amount of hydrogen sulphide was determined by iodometric titration. As one molecule of H,O is used for the formation of each molecule of H,S, the equilibrium ratio H,S/H,O would be , where (H,O) is the molar concentration in the ingoing gas, and (H,S) the molar concentration in the outgoing gas. In the present work (H,S) was always very small compared to (H20). In order for the observed H,S/H20 ratio to represent the true equilibrium ratio the gas flow has to be: 1—Sufficiently slow to give a complete establishment of equilibrium, and 2—sufficiently fast to counteract thermal diffusion. Incomplete reaction would give a value decreasing with increasing flow rate, and thermal diffusion would give a value increasing with decreasing flow rate. When inlet and outlet tubes of about 2 sq mm cross-section were used, the observed gas ratio was independent of the flow rate between 15 and 125 cc per min, Fig. 2. In this range, therefore, the observed gas ratio represents true equilibrium.* For the rest of the in-

Jan 1, 1952

By Raynor Linzey, Karl M. Busen

The formation of thin films of silicon oxide by anodic oxidation of silicon and the subsequent removal of these films by an etch is a process often used for the evaluation of concentration distributions Profiles) in silicon layers by the differential sheet conductance method. The accuracy of the resulting profile is very strongly influenced by the uniformity of the thickness of the reacted silicon. Normally, it would be expected that for a constant number of coulombs passed the thickness would be the same from oxidation to oxidation. Investigations show that in certain electrolytes, for a given number of coulombs passed through an n-type silicon sample, the thickness of the reacted silicon increased with increasing resistivity. Even for the same resistivity the thickness varied sometimes by a factor of 1.5. When an electrolyte was used which consisted of 10 pct water by volume in ethylene glycol with 4.0 g KNO per 1000 ml of solution, anodiza-tion at 5 ma per sq cm led to satisfying results. Short-time anodizations gave oxide layers of a higher apparent density than those experienced from thermally gown silicon oxides. THE functioning and the electrical characteristics of semiconductor devices are based upon the incorporation of "impurities" into a single-crystalline body of suitable material and on the concentration distribution (profile) of this impurity. The incorporation can be achieved by well-known processes as, for example, by diffusion, epitaxial growth, alloying, or ion implantation. Often the profile resulting from these processes is not known. A powerful tool to learn about a concentration distribution is given by the method of differential sheet conductance which employs successive four-probe measurements on a layer subjected to stepwise removal of thin sublayers. Differential sheet conductance vs position (or penetration depth) of a sublayer is plotted and a smooth curve is drawn through the data. From this curve the profile is then calculated. When the semiconductor material is silicon, the sublayers are removed suitably by anodic oxidation of the silicon and subsequent dissolving of the formed silicon oxide by an etch. The accuracy of the resulting profile is very strongly influenced by the uniformity of the sublayer thickness. Normally, it would be expected that, for a constant number of coulombs passed, the thickness would be constant from oxidation to oxidation. Investigations showed that for sublayers several hundred angstroms thick the reproducibility can be rather poor. Therefore, efforts were made to obtain a reliable technique for uniform removal. The present paper describes such a technique and certain phenomena which were encountered during the investigations. APPARATUS AND TECHNIQUE The first report on the investigation of concentration distributions, where for differential sheet conductance measurements thin silicon sublayers were removed from a diffused layer by anodic oxidation, was given by Tannenbaum. The author reports the removal of sublayers which for the most part were 400 thick. More advanced device designs now require much narrower layers. When it was tried in these laboratories to determine profiles within such layers, difficulties were encountered with respect to sublayers which by necessity had to be thinner than the ones reported by Tannenbaum. The apparatus used for the investigations is sketched in Fig. 1. Two cylindrical containers connected by a wide tube are filled with an electrolyte. The left container receives the cathode whereas the right container is closed at the bottom end by a sample support consisting of a Teflon base and a copper pedestal. The silicon sample (1 by 1 cm) is mounted to the pedestal embedded flush in the Teflon base using silver paint (Degussa) for electrical contact. Pyseal is applied to the edges of the sample to protect the copper pedestal from the electrolyte. The base is mounted to the container using a water-tight silicon rubber gasket. Fig. 2 gives a view of the sample support. The copper pedestal is connected to the positive side of a power supply operating at a constant current output (Kepco Model ABC 425M). The electrolyte which has been reported by Duffek et al. consisted of either 2 or 10 pct water by volume in ethylene glycol and 4.0 g KNO3 in 1000-ml solution. The electrolyte is best prepared

Jan 1, 1967

By M. Cohen, D. J. Blickwede

Quantitative observations concerning the carbide phases in high speed steel are of importance for two general reasons: (1) the carbides, being inevitable constituents of the final structure, exert a direct influence on the properties of the steel; and (2) a substantial proportion of the total alloy content is tied-up in the carbides, and hence the extent of their solution on austenitizing governs the composition of the steel matrix. The latter relationship has a vital bearing on the response of the steel to tempering as well as on its performance in subsequent service. Accordingly, in the course of a long-term study of the behavior of high speed steels, the authors were confronted with the problem of securing quantitative data on the carbide phases. The obvious method for acquiring such information is to isolate the car-bides from the steel and subject them to chemical, X ray diffraction and other measurements. There are well-known extraction techniques which involve the chemical or electrolytic solution of the less noble matrix (ferrite, marten-site or austenite), thus leaving a residue of the carbide phases. However, the results obtained must be scrutinized carefully1,2 since the carbides may be affected by the chemical or electrolytic action. It is the purpose of the present paper to describe the experiments leading to an electrolytic-extraction technique for quantitatively isolating the carbides from both I and hardened high speed steel. Particular attention is paid to the amount, as well as the composition, of the carbides so that the matrix analysis becomes ascertainable by subtraction from the overall steel composition. Illustrative results are given for the M-2 grade of tungsten-molybdenum steel. Review of the Literature The chemical method of dissolving the matrix selectively with respect to the carbides makes use of dilute non- oxidizing reagents such as hydrochloric or sulphuric acid. Although this simple procedure has led to the determination of the cementite composition,3,4 it achieved only limited success because of the interaction between the acid and the carbide residue. Some of the carbides may not only be destroyed in this way, but the hydrogen released is likely to remove part of their carbon as hydrocarbon gases. The electrolytic technique of isolating carbides has the advantage of rapidly dissolving the specimen (anode) in the presence of less reactive solutions than are practicable with the chemical method. This reduces the possibility of chemical attack on the carbides, and furthermore, the hydrogen evolved during the electrolysis is released at the cathode which is not in close proximity to the carbides. The common electrolytes adopted for this purpose are hydrochloric and sulphuric acids.5-l1 Aqueous solutions of ferrous salts have also been used.12,13 A considerable advance in experimental technique was introduced by Treje and Benedicks14 who developed a double-compartment cell for electrolytic extraction, the anode and cathode chambers being separated by a porous diaphragm. A solution of 15 pct sodium citrate, 2 pct potassium bromide and 1 pct potassium iodide was selected for the anolyte, while the catholyte consisted of a 10 pct solution of copper sulphate, with copper serving as the cathode. This type of cell has a number of desirable characteristics: 1. The anolyte has a pH value close to 7, at least at the beginning of the run. 2. The iron that dissolves from the anode-specimen forms a water-soluble complex ion with the citrate, thereby preventing the precipitation of iron hydroxide (which would contaminate the carbide residue) despite the neutrality of the solution. 3. Copper deposition instead of hydrogen evolution occurs at the cathode, and this avoids an increasing concentration of hy-droxyl ions which (in an otherwise neutral solution) might cause the precipitation of insoluble hydroxides. 4. Contamination of the anode chamber by copper sulphate is inhibited by the porous diaphragm. Houdremont and coworkers15 applied the above method (with the further refinement of excluding oxygen during the electrolysis, washing and drying) to the extraction of carbides from a series of plain carbon steels after various heat treatments. They had quantitative success only with specimens in the annealed condition, and concluded that the size and shape of the carbide particles play an important role in the isolation process, with large spheroids exhibiting the least tendency to decompose during the electrolysis. Up to the present time, the citrate double-cell has not been used to any extent for isolating the carbides of high alloy steels, apparently on the grounds that the complex carbides are more resistant than cementite to attack in the simpler acid electrolytes. In particular, Bain and Grossmann7 and Gulyaev10.10a have employed the hydrochloric acid cell for their investigations of the carbides in high speed steel.* It will be demonstrated here that this type of cell is capable of yielding quantitative results in the case of high speed steel, and actually has certain advantages over the more complicated double cell. However, in order to provide a rigorous test of the quanti-tativeness of electrolytic procedures for the problem at hand, both methods were studied in considerable detail.

Jan 1, 1950

Mining as an industry of Utah had its inception in the activities of United States soldiers who came to the Salt Lake valley under the command of Gen. P. E. Connor, founder of Fort Douglas, in October, 1862. Many of these soldiers were from California and naturally employed their spare time in the search for minerals. Although their military status did not permit them to make locations in their own names, they assiduously prospected the region in accordance with their own knowledge and with in- formation conveyed to them by Indians and others. Their efforts were encouraged by their commander, who thereby fairly earned the title of father of the industry here, and by their other officers, and many discoveries that subsequently proved of great value were made by the soldiers. No systematic prospecting or mining had been done in the Territory previously although early Mormon settlers in the southern portion had mined a little ore from time to time in order to provide supplies of lead for bullets. The Mormon leaders at that time did not encourage mining. They held that agriculture was more important and profitable for their followers.

Jan 1, 1925

By F. Forscher

Pronounced porosity, decreasing with distance from the fracture surface, is found in the necked region of tensile specimens tested at room temperature or liquid nitrogen temperature. A hydrogen solution treatment followed by a quench prevents pore formation in tests at —196°C, but has little effect on tests made at room temperature. Reduction of area values at —196°C are strikingly improved by the same treatment, but show no improvement in room temperature tests. The experimental results support the hypothesis that strain-induced porosity results with hydride initially present or, in material containing less than 50 ppm H, with hydride precipitated by strain aging. ZIRCONIUM'S advent as an engineering material in nuclear reactor construction has brought forth many investigations concerned with the properties of this metal and its alloys. Some attention has been given to the ductility of the metal.1, 2 Particularly, the effect of hydrogen on the impact strength was noted," indicating a pronounced dependence of ductility on the hydrogen content and its distribution in the metal. Micro graphic:' and X-ray diffraction studies' have shown that a precipitated hydride phase causes embrittlement in zirconium,' which implies a distinctly different mechanism from that causing hydrogen embrittlement in steels."-' In the latter case, it is assumed that atomic hydrogen diffuses to rifts, where it forms molecular hydrogen under such internal pressures that it causes the rifts to enlarge, leading to fracture. Thermo-dynamic studies of the Zr-H system8, 9 also indicate that the internal pressure of molecular hydrogen would be too low to cause any damage to the zirconium at room temperature, and the formation of a zirconium hydride (d phase) is favored." It is noteworthy that this phase has a face-centered-cubic structure and nearly 18 pet larger molar volume than close-packed-hexagonal a-zlrconlum' Other features of the ductility of zirconium are the absence of a typical cup-cone failure upon tensile testing, a pronounced strain-induced porosity in the necked region of a tensile bar, Fig 1, and the associated spongy appearance of the fracture surface A tremendous improvement in ductility at liquid nitrogen temperature can be observed following a heat treatment that retains the commonly present quan- tities (10 to 50 ppm) of hydrogen in solution. Such tests show a typical cup-cone failure and the absence of strain-induced porosity, Fig. 2. The present paper reports results of an investigation into the causes of strain-induced porosity in zirconium. Room temperature tensile fractures of mild steel, aluminum, and titanium have also been examined for strain-induced porosity. In each case porosity could be detected, but the amount was orders of magnitude less than was found in zirconium. Of the three metals, titanium showed the strongest effect. Material, Specimens, and Preparation The greater part of this investigation was done on Westinghouse crystal bar zirconium, are-melted, forged, and hot-rolled. Chemical analyses of the hot-rolled plates are given in Table I, along with those of the Foote crystal bar zirconium used in this work. The specimens were machined from the hot-rolled plate (1550°F) of 1/2 in. thickness and subsequently heat-treated according to one of the treatments shown in Table IS. All specimens were machined round, the longitudinal axis in the rolling direction, with a gage

Jan 1, 1957