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By J. B. Clark

An exploratory study of the 60 wt pct Mg and 80 wt pct Mg sections of the solid constitution in the magnesium-rich region of the Mg-Al-Ca-Zn system was undertaken. The objective of the investigation was to determine the general configuration of the phase fields with respect to the magnesium solid solution in this complex system. The 335°C (635°F) tetrahedron was selected for study because the solid constitutions of the bounding ternary systems—Mg-Al-Zn,1 Mg-A1-Ca,2 Mg-Ca-zn3—were well known and, at least in these ternary systems, diffusion was rapid enough at 335°C for equilibrium to be reached fairly rapidly. Metallography and powder X-ray diffraction were selected as the best corroboratory methods for determining the solid phase equilibria. A total of forty-six quaternary alloys on the 60 wt Mg and 80 wt Mg sections were prepared, annealed to equilibrium at 335°C, and the phases present identified. Singly sublimed (99.99 pct) magnesium, electrolytic (99.99 pct) zinc, high-purity (99.99 pct) aluminum, and USP calcium were alloyed in graphite crucibles under a chloride flux and chill cast into a split graphite mold. Alloy samples were sealed under argon in Pyrex vials and annealed at 335°C ± 1/2° for a minimum of 500 hr. At the conclusion of the anneal, the alloy samples were quenched rapidly in ice water. Standard metallographic procedures for magnesium alloys were used. Powder X-ray diffraction patterns were taken on a G. E. Camera with a diameter of 14.32 cm using nickel filtered copper radiation. The phase equilibria in the 60 wt pct Mg section of the 335°C (635°F) tetrahedron is shown in Fig. 1. The complex equilibria in the zinc-rich region were not accurately determined and are shown only tentatively with dotted lines. The equilibria of the three bounding ternary 335°C sections from pure magnesium to 60 wt pct Mg are also shown as though folded away from the tetrahedron. The 80 wt pct Mg section has the same configuration of phase fields as that shown in Fig. 1. Thus it is reasonable to conclude that in higher magnesium sections, the same relative configuration of phase fields is maintained though the relative proportions of the individual phase fields may change slightly as the magnesium solid solution is approached. Note that all the ternary phases present in these ternary diagrams enter into the quaternary equilibria with the exception of the phase of the Mg-Al-Zn system which becomes unstable with the addition of as little as 1/2 pct Ca. No quaternary intermetallic phases were found. No liquid phase was detected in any of the quaternary alloys annealed at 335°C suggesting that if a quaternary eu-tectic exists, it must lie between 335°C and 339°C, the temperature of the ternary eutectic in the Mg-Al-Zn system.

Jan 1, 1962

By John H. Hall

The Frémont test for measuring the energy consumed in breaking a notched bar of steel is not so well known in this country as it deserves to be. The test specimen used in this test is about # by 6 by about 16 in., notched with a hack saw on the lower side to cause the piece to break at a fixed point. A weight of 10 kg. falling 4 m., and provided with a knife or chisel which strikes the specimen, breaks the test piece by impact. The residual force in the falling weight after it has broken the specimen is recorded by a suitable device, and the difference between the total energy and the residual energy of the weight (in kilogrammeters) is the force consumed in breaking the specimen. An experience of several years with the Frémont testing machine has shown its great value in readily detecting brittleness not revealed by tensile testing or slow bending testing. In this paper it is proposed to discuss (1) the supersensitiveness of the test as applied to cast steel and (2) its value in revealing the effect of heat treatment under certain conditions. In testing coupons cut from castings heat treated by quenching and annealing, attention was drawn to the wide variations in the value of the Frémont test of specimens cut from different parts of the coupon. The latter was of the form and dimensions shown in Fig. 1, and Fremont test pieces were cut from two parts of this coupon, marked A and B in that figure. The results of the tests are shown below, tests A and B in each instance being from the same coupon: A. B. A. B. Coupon. Kilogram- Kilogram- Coupon. Kilogram- Kilogram- meters. meters. meters. meters. 1,........ 12.0 12.5 9,........ 27.5 32.5 2,........ 11.0 32.5 10,........17.5 31.0 3,........ 11.5 30.0 11,........ 14.5 31.0 4,........ 20.0 33.0 12,........ 15.0 27.5 5,........ 17.5 33.0 13,........ 19.0 30.0 6,........ 17.5 20.0 14,........ 17.5 25.5 7,........ 17.5 13.0 15,........15.0 30.0 8,........ 10.0 30.0

Jan 1, 1914

By J. M. Woodall

GuAs crystals which have been grown in quartz boats by the horizontal Bridgman method in the pvesence of Ga20 vapov have beetz found to have carrier and donor distributions which do not correspound to those expected from simple dopant seg-vegation during directional freezing; Instead, the carrier distribution is determined by the heat-trentnzent history of the crystal, while the donor distribution, zohiclz is principally due to silicon, is fixed by the pozuth rate, the geo)tlet.ry of the crystal growth vessel, and tlze initial Ga20 pressure. WHEN semiconducting materials are made into doped crystals by the normal freezing method,&apos; they usually exhibit doping variations along the growth axis. If a) there is no dopant diffusion in the solid, b) the dopant is distributed uniformly in the melt, and c) the distribution coefficient, k, does not vary with composition, then the doping variation along the growth axis is represented by the equation: where C is the doping level in the solid at a point where a fraction g of the original liquid has frozen, and Co is the mean concentration. When the dopant is either a singly ionized shallow donor or shallow acceptor, C also represents the carrier concentration. Even though this equation accurately describes the dcping profiles of a large number of normal freeze systems, there are several special systems for which Eq. [I] does not apply. One such system is the horizontal Bridgman method for preparing oxygen-grown GaAs crystals using quartz vessels. Several workers2"* have shown that GaAs crystals grown by the horizontal Bridglnan method using quartz vessels are generally contaminated with silicon in concentrations in excess of 5 < 10lG atoms per cu cm. This contamination is ascribed to a reaction: occurring at the walls of the crystal-growth vessel which liberates silicon into the melt and Ga2O vapor. It has been shown4 that this reaction, and, hence, the silicon contamination, can be suppressed by the addition of oxygen to the crystal-growth apparatus. It is the purpose of this paper to describe a special apparatus capable of yielding single crystals of GaAs grown in the presence of oxygen and to describe both the kinetics of silicon suppression in this system and the relationship between the carrier concentration profile and the silicon concentration profile. EXPERIMENTAL-CRYSTAL GROWTH A schematic diagram of the apparatus used in the oxygen addition experiments is shown in Fig. 1. The most important features of this apparatus are: a) the use of a sand blasted quartz boat, h) a quartz rod of length 1, with a hole of cross section A, that is placed near the boat to limit the free space volume V, over the melt during the growth, and c) the temperature gradient at and near the solid-liquid interface. Sand blasting the boat is necessary to prevent wetting of the melt. The quartz rod retards the diffusion of the GazO vapor away from the melt to the colder portions of the ampoule. A crystal is prepared by first loading a 5.5-in. boat, which has been cleaned in aqua regia, with 40 g of 99.9999 pct Ga along with 1 to 8 mg of Ga203 powder. GazO3 is a convenient source of oxygen since it reacts with gallium at the melt temperature to form Ga20 vapor, the species which apparently controls the suppression of silicon contamination. The loaded boat, the quartz rod, and the 99.9999 pct As are placed in the ampoule as shown in Fig. 1. Generally, GaAs seeds were not used since most of the unseeded growths resulted in monocrys-tals. The ampoule is evacuated to 10-5 Torr and sealed off. The GaAs melt is synthesized by placing the ampoule into a two-zone horizontal Bridgman furnace. The two zones are separated by several sandwiched layers of -in. Fibre-Fax board drilled with holes slightly larger than the OD of the ampoule. This causes a very large temperature gradient between the two zones, which is necessary for single-crystal growth. The two-zone furnace is mounted on a stand fitted with a roller bearing which allows uniform motion of the furnace in a horizontal direction. A uniform velocity of the stand is achieved by the use of a windlass device which winds up a wire attached to the stand. Movement of the solid-liquid interface is accomplished by fixing the position of the ampoule and moving the stand. Growth rates investigated in this experiment were between 0.4 and 1.2 in. per hr. The composition of the melt is fixed by maintaining an arsenic reservoir at 618°C. The stand is moved away

Jan 1, 1968

By C. M. Sellars, F. Maak

Electvomotie -force measurements hazle been made on ten Au-Ni alloys at temperatures 7754 825O, 900O, and 935°C using galvanic cells with solid electrolyte. Partial and ivtegral thermodynamic functions are derived from the determined activities. which show la9,ge positive deviations jvom Raoult&apos;s law. The enthalpies of mixing- are in close accord with calorinletric data and are of comparable accuracy. The results are used to determine the miscibility gap and the spinoidal curL1e, and are qlalitaticely interpreted in terms of the large size difference and electvonic interactions between gold and nickel. Tabulated nalues of the actielities and partial and inteqal free energies, entropies , and enthalpies of mixing are given in the appendix. The Au-Ni binary alloys are of particular interest thermodynamically as despite the large size difference between gold and nickel atoms they have a simple phase diagram.&apos; This shows complete solid solubility above a miscibility gap with a maximum at about 71 at. pct Ni and 810°C. Activity measurements (over a range of temperature) have been made previously on this system using an electrolytic-cell method with a liquid chloride eletrolyte. The accuracy of this data has, however, been queried3&apos;4 as it leads to values for the enthalpy of mixing which are too high when compared with calorimetric measurements.46 Also, the entropy of mixing is too high when compared with the values derived from specific-heat measrements. It was therefore considered worthwhile redetermining the activity data using a galvanic cell with a solid electrolyte (0.85 ZrO, 0.15 CaO). This method is based on a technique introduced by Kiukkola and aner&apos;&apos; and has been used to determine the thermodynamic data for Cu-Ni alloys.&apos; THE ELECTROLYTIC CELL The cell used to determine the activity of nickel in Au-Ni alloys was: The cell reaction, which is discussed for a similar cell by Rapp and Maak, gives a direct measurement of the activity of nickel in the alloys as where aNi(alloy) is the activity of nickel in the alloy, 3 is Faraday&apos;s constant, E is the measured electromotive force, R is the gas constant, and T is absolute temperature. For the valid application of this type of cell the arrangement must be ther modynamic ally stable, which means that: 1) The free energies of formation of the oxides fulfill the relation 2) The vapor pressures or dissociation pressures of each component must be so small that during the time of the experiment the amount of any material transported by the gas is negligible. 3) Side reactions which change the alloy composition must be negligible. These reactions are possible both with the surrounding gas atmosphere and with neighboring solid material. The first two conditions are fulfilled by the above cell.&apos; The third condition is approached by care in the experimental technique. The factors which control the attainment of equilibrium in this type of cell have been considered by Rapp and aak, who conclude that metallic diffusion in the alloy is the rate-control ling step. A lower temperature limit for the practical application of this cell is therefore set by the condition that equilibrium must be attained in reasonable times. EXPERIMENTAL MATERIALS AND TECHNIQUE The starting materials were mint gold (99.99 pct Au) and Mond nickel (99.9 pct Ni). Alloys, nominally 5 at. pct Ni and then each 10 at. pct across the system, were produced by vacuum-levitation melting compacts of mixed drillings of the pure metals. The nickel used for the reference tablets in the cells was melted in a similar manner. The alloy ingots, cast in boron nitride molds, were cold-forged and

Jan 1, 1967

By Francis Lucas

EACH year we gather in this auditorium to honor the memory of a . distinguished American metallurgist and educator. I cannot bring to you reminiscences of Prof. Henry Marion Howe as other lecturers have done. It was not my good fortune to enjoy the privilege of studying and ,working under his direction. As I have reflected on the published works of Professor Howe, it has seemed to me that his mind was trained on the nature of metals and on the mechanism which produced known but unexplained effects. He was interested in controlling and regulating the causes so as to produce specific effects, useful in serving mankind. I recall an afternoon in May, 1925, when, at the invitation of Mrs. Howe, I visited his former home, Green Peace at Bedford Hills, N. Y. The purpose of my visit was to receive at the hands of the widow -an engraving of Professor Howe, a reproduction of a painting from life. Mrs. Howe showed me his study. The walls of this large attractive room, brilliantly illuminated by the spring sunshine, were literally covered with the scrolls of honors and of decorations which had been conferred on Professor Howe by governments, academies, universities, societies and. institutions the world over. As we walked about the room, she pointed out the significance of each honor and recited a few details of each occasion. I sensed the pride with which Mrs. Howe, then in the late autumn of her life, told me of her husbands&apos; achievements, of his struggles, of his disappointments and of his successes. There was a surprisingly familiar note in the recital to which I was listening. I have found the same theme running as a disconnected thread through the written accounts of the lives of most men of science. We then went to the drawing room to chat for a few minutes before I said my farewells. As we seated ourselves, Mrs. Howe turned to me and asked: "Mr. Lucas, what of this great microscope of yours? Cannot it be applied to the study of disease? What a great blessing it would be if it could be directed to relieve the suffering of mankind."

Jan 1, 1931

By R. L. Fullman

IN the investigation of metallurgical transformations and the relationships between microstructure and properties of metals, it frequently is desirable to obtain a measurement of the relative amounts of the various phases present and of the mean size of particles into which each phase is dispersed. The relative amounts of the phases can be measured by the classical methods of area, lineal, and point analysis,1-5 in accordance with the principle that the volume fraction of a phase, the fraction of a polished cross section occupied by the phase, the fraction of a random line occupied by the phase, and the fraction of randomly arrayed points occupied by the phase are all equal. The validity of this relationship depends only on the attainment of a truly random sample of area, length, or points, and not on the size, shape, or distribution of the particles constituting the phase. Smith and Guttman8 have derived a relationship between the interface area per unit volume S, and the measurable quantities L., the interface length per unit area on a cross section, and NL, the number of interfaces per unit length intersected by a random line. Their equation, Sv = — L8 = 2NL is also valid regardless of the distribution of particle sizes and shapes. In contrast to the situation concerning measurement of relative fractions of phases and of interface area, the measurement of particle sizes in opaque samples has not been subjected to a complete analysis. It has been common to measure some lineal or area dimension of particles on a polished cross section and to use the mean value as a qualitative measure of particle size. In the present paper, quantitative relationships are established among the various mean dimensions on a polished cross section and the actual dimensions of the particles present. Particles of Uniform Size Spheres: If a metal sample contains particles of a phase a dispersed in the form of spheres of uniform size, a polished cross section through the sample will reveal circular areas of phase a with radii from 0 to ?, the radius of the spheres. Consider a cube of unit dimensions to be cut from the sample. If a cross section parallel to one of the cube faces is examined, the average number of particles per unit area (N,) equals the number of particles per unit volume (Nv) times the probability p1 that the plane would intersect a single sphere positioned at random within the unit cube. Since, of the various possible positions for the cross-sectional plane over the unit length from top to bottom of the cube, only those positions existing over the length 2r would lead to the plane intersecting the sphere, the probability of intersecting a single sphere is just 2r. N8= Nvp1 = Nd-2r [1] Applying the equality of area and volume fractions, the relationship is found between sphere size and average area s of uniform spheres intersected by a random cross section, 4 - f = NV V = Nr . — pra = N s = Nd . 2rs S = —pr2 [2] A similar analysis reveals the average traverse length across spheres of uniform size when random lines are passed through the sample. If a randomly oriented unit cube is cut from the sample and a randomly positioned line is passed through the cube parallel to a cube edge, the number of spheres intersected by the line (Nl) equals the number of spheres per unit volume times the probability p1 of the line hitting a single randomly placed sphere in the cube. Since possible positions of the line occupy unit area, and possible positions for which it will pass through the sphere occupy an area of pr2, the probability of the line hitting a randomly placed single sphere is pr2. NL = Nv p1 = Nvpr2 [3] Combining this relationship with the equality of volume and lineal fraction, the desired relationship is obtained between radius and mean lineal traverse length -i, for spheres of uniform size. 4 - - 3 l=4/3r [4] Circular Plates: Consider a sample containing particles of a phase a in the form of circular plates of uniform radius r and thickness t, where r >> t. If the plates are randomly oriented, as in a sufficiently large sample of a fine grained polycrystalline material, area and lineal analysis may be carried out with parallel cross-sectional planes and lineal traverses. If the plates are not randomly oriented, it is necessary to randomize the orientation of the cross-sectional planes and traverse directions. Let a unit cube be cut from the sample, and a cross-section plane be passed through the cube parallel to one of the cube faces. The number of plates cut by the cross-sectional plane per unit area is equal to the number of plates per unit volume times the probability of a plate intersecting a single randomly positioned and randomly oriented plate in the cube. If J is the component of the plate diameter in the direction normal to the cross-sectional planes, the probability of a plane cutting a single randomly oriented plate is equal to J, the mean value of J for all possible orientations of the plate. Let 4 be the dihedral angle between a plate and the cross-sectional plane, and let p?, d? be the probability that a plate makes an angle between 4 and ? + d? with the cross-sectional plane. Then for ran-

Jan 1, 1954

By M. J. Marcinkowski, A. Szirmae, R. M. Fisher

Room -temperature hardness measurements obtained from single and polycrystalline samples of a 47.8 at, pet Cr-Fe alloy which were aged for various times al 500°C show a two-fold increase over that of the unaged alloy after annealing for 1000 hr. A detailed examination of the deformation markings in the neighborhood of the hardness imbressions reveals that twinning becomes an increasingly more important mode of deformation as aging proceeds. this observation is shown to be inconsistent with the Proposals that the transformation is eve of order-disorder On the other hand, the results are in agreement with previous observations of Fisher et al. of a coherent chromium-rich precipitate which forms in an iron-rich matrix as a result of a miscibility gap in this alloy system as proposed by Williams . Using the coherent precipitate hypotlzesis as a model, a detailed analysis of the various possible strengthening contributions to both the slip mid twinning stresses is made. In the case of&apos; slip, lattice friction within the chromium-rich phase and the chemical energy associated with the interface between the two phases contribute about 60pet to the total strength. The contributions from coherency strains and modulus differences are thought to contribute the remaining 40 pet of the strength but are difficult to evaluate because of uncertainties regarding the flexibility of the dislocation line. All of the factors have nearly the same or else a smaller effect on the twinning stress in the aged alloy. Twinning is never observed in the unaged alloy because the twinning stress is much higher- than that for slip. With increased aging times, the various contributions to the total stress for both twinning and slip increase, but most of those for slip increase much more rapidly, so that, in the fully aged alloy, it surpasses the stress for twin propagation. When a twin is nucleated by the chance occurrence of an internal stress concentration during a test, the subsequent twin burst results in a low hardness reading. FeRNTIC Fe-Cr alloys from about 15 to 80 at. pet Cr content show considerable change in properties such as hardness, electrical resistivity, saturation magnetization, and so forth, after aging in the vicinity of 500°C.* The most marked change is a large increase in hardness accompanied by a sharp decrease in ductility, so the phenomenon has often been referred to as the 475°C (885°F) em-brittlement problem. Changes in microstructure are rather subtle and most investigators have not been able to observe any X-ray diffraction or metal-lographic changes during aging. There is little doubt that the phenomenon is inherent to the Fe-Cr binary system and is not directly related to the formation of a phase. Two distinct schools of thought have developed during the past decade concerning this problem. Masumoto, Saito, and sugiharal have concluded, from their own specific-heat measurements, that atomic ordering occurs in this temperature range. Pomey and Bastien2 have also attributed the changes in physical properties with aging in the neighborhood of 500°C to atomic ordering. In addition, Takeda and Nagai3 claimed to have found X-ray verification for superlattices corresponding to the compositions FesCr, FeCr, and FeCr3. However, all known attempts to observe superlattices in Fe-Cr alloys using neutron diffraction, which should be much more sensitive, have been unsuccessful. These have been reported by Shull et al.4 for 25 at. pet Cr, williams5 for 75 at. pet Cr, and Tisinai and samans6 for a 28.5 at. pet Cr. steel. The second and more prevalent opinion ascribes the 475°C embrittlement phenomenon to the formation of a coherent precipitate due to the occurrence of a miscibility gap in the Fe-Cr system below about 550°C. This latter concept was first indicated by the results of Fisher, Dulis, and Carroll,7 who were able to extract fine particles about 200Å in diameter from samples of 28.5 at. pet Cr steels aged from 1 to 3 years at 475°C. The extracted material was found to be nonmagnetic, to have a bee structure with a lattice parameter between that of iron and chromium, and to contain about 80 at. pet Cr. Williams and paxton8 and williams5 confirmed these results and first explicitly proposed the existence of a miscibility gap below the a region in the equilibrium diagram. Alloys aged within this gap

Jan 1, 1964

By C. Atkinson

A method is described for the numerical solution of the diffusion equation with a composition-dependent diffusion coefficient and applied to the radial growth of a cylinder; the radial growth of a sphere, and the symmetric growth of an ellipsoid. Sample applications of the method are made to the growth of particles of proeutectoid ferrite into austenite. RECENTLY&apos; we described a method for numerical solution of the diffusion equation with a composition-dependent diffusion coefficient for the case of the growth of a planar interface. In this paper we extend this method to describe the radial growth of a cylinder, the radial growth of a sphere, and the symmetric growth of an ellipsoid. In the latter case, limiting values of the axial ratios of the ellipsoid reduces the problem to one of a cylinder, a sphere, or a plane depending on the axial ratio. A check on these limiting values is made in the results section. In all of these cases we consider growth from zero size. A natural consequence of this assumption as applied to the sphere, for example, is that the radius of the sphere is proportional to the square root of the time. This is consistent with the condition that the radius is zero initially, i.e., grows from zero size. It may be argued that it is more realistic to consider particles which grow from a nucleus of finite initial size; even in this case the analysis of this paper is likely to be applicable. This can be seen if a comparison is made of the work of Cable and Evans,2 who consider a sphere of initially finite size growing by diffusion in a matrix with a constant diffusion coefficient, with the results of Scriven3 for growth from zero size. This comparison shows that the rates of growth in each case differ trivially by the time the particle has grown to about five times its initial size." This investigation is a generalization of those of Zener,4 Ham,5 and Horvay and cahn6 to the situation often encountered experimentally, in which the diffusion coefficient varies with concentration. First let us consider each of the cases separately. I) GROWTH OF SPHERICAL PARTICLES FROM ZERO SIZE In this case the differential equation in the matrix depends only on R, the radius in spherical coordinates, and can be written: ? 1 <^\ ^13D . , dt U\dRz + R 3Rj + dR dR [ J where C is the composition, t is the time, and D is the diffusion coefficient which depends on c. The boundary conditions will be: c = c, at the moving interface in the matrix, c = c, at infinity in the matrix (and at t = 0, everywhere in the matrix), c = X, is the composition in the spherical particle. Each of the above compositions is assumed constant. In addition there is the flu condition at the moving interface which can be written: , dR0 ~/3c dt \dR/H =Ra where R,, which is a function of t, is the position of the moving interface. We make the substitution q = RI~ in [I] reducing this equation to: & - m - *ws) »i where we have written D = D,F(c) or simply D,F, and Do = D(c,). Thus F[c(q0)] = 1 where q, = ~,/a is the value of the dimensionless parameter q evaluated at the interface. Multiplying Eq. [2] by dq/dc and integrating, we find: where the lower limit of the integral has been chosen so that dc/dq — 0 as c — c,, thereby satisfying the boundary condition at infinity. We require, then, to solve Eq. [3] subject to the condition c = c, when q = q, (this follows from putting R = R, at the interface) together with the flux condition which can be rewritten in terms of q as: Eqs. [3] and [4] together with the condition c = c, at q = q0 enable us to find 77, and the concentration profile c = c(q). Numerical Method. We treat Eq. [3] in the same way as we did the corresponding equation for the planar interface problem&apos; i.e., by dividing the interval c, to c, into n equal steps so that: cr = ca -rbc [5] where r takes the values 0, 1, ... n and we call no,, q1, ... nn the values of n corresponding to the compositions c,, c,, ... c,.

Jan 1, 1970

By John V. Beall

There is an underlying surge to mining in the Philippines." So said A. R. "Ham" Reed, general manager Itogon-Suyoc Mines, Inc., to this reporter last November. Itogon-Suyoc is a small company with a struggling gold mine, a mixed sulfide mine and some prospects. In the Republic of Philippines (ROP), politics, which have an effect on mining, create some uncertainties. By July 1974, no mining company may have more than 40% foreign ownership. The ownership provision is a matter of concern to such important companies as Atlas and Benguet Consolidated and is generally considered a deterrent to new foreign investment. Under these circumstances, Reed&apos;s comments are of interest. On foreign prospectors: Don&apos;t come in cold; find a Philippine partner or buy into a company already here. Physical conditions of exploration are tough-terrain, communications -but the mineral potential is excellent.

Jan 4, 1969

By G. R. Fitterer

NONMETALLIC inclusions in steel have received much attention by metallurgists during the past few years. Many investigators have been emphatic in stating their belief that these impurities are the chief causes of steel failures, while others think that their effects are negligible. One of the main reasons that definite information on the effects of these impurities is lacking is that no reliable method has been developed for the exact determination of the amounts of these different impurities in steel. Some valuable new information has been brought out by a reliable method for nonmetallic inclusions which has been developed at the Pittsburgh Experiment Station of the U. S. Bureau of Mines, in cooperation with the Carnegie Institute of Technology and the Metallurgical Advisory Board. This procedure satisfactorily extracts some of the more common inclusions from steels, thus enabling the investigator to correlate steel failures with inclusion content. Although the data given herein are the first to be published on the subject of the electrolytic extraction method, the interest in the subject is evidenced by the fact that it is already being used by some 20 to 30 steel companies for their various purposes. Some of the comments received from these laboratories indicate that the method is highly satisfactory: the unfavorable comments, when traced to their sources, have been found to be based on some small mistake made in either installation or operation. The chief discrepancy seems to have occurred in attempts to apply the method to the extraction of inclusions from alloy steels. The procedure can be used only for the plain carbon steels and for such low alloy steels as those containing the usual amount of copper (0.2 to 0.3 per cent). Steels containing the carbide-forming alloys such as chromium, tungsten, etc., present far too many difficulties for the accurate determination of inclusions. Luckily, the iron and manganese carbides do not offer any difficulties and no contamination of the oxide-sulfide residue results. It is highly probable that some modification of the following

Jan 1, 1931

By E. S. Machlin, Morris Cohen

The isothermal formation of martensite in a 71 pct Fe, 29 pct Ni alloy is found to take place mainly by the nucleation of new plates rather than by the growth of existing ones, and is dependent on the temperature and time of the isothermal hold, the amount of atherrnal martensite present, the state of internal strain, and the application of tensile stress. The conclusion is reached that the isothermal nucleation is activated by thermal fluctuations superimposed on localized regions of very high strain. The reaction-path concept, involving martensite nucleation at strain embryos by athermal activation on cooling, or by thermal fluctuations on holding, reconciles the athermal and isothermal modes of transformation, and offers a unified picture of the kinetics of martensite formation. THE phenomenon of isothermal martensite formation has been the subject of considerable study lately,&apos;-&apos; and its reality has been thoroughly established. The existence of such isothermal transformation has been claimed to confirm the nucleation and growth hypothesis as applied to the martensitic reaction.1 " However, the isothermal formation of a martensitic phase is also compatible with the reaction-pathw escription of the transformation. The purpose of this investigation is to examine the isothermal mode of the martensitic reaction, both theoretically and experimentally, in order to test the likelihood of the above two mechanisms. Material and Methods The alloy used in this study contained about 71 pct Fe and 29 pct Ni. The complete analysis was 29.5 ± 0.2 pct Ni, 0.036 pct C, 0.02 pct N, 0.19 pct Mn, 0.09 pct Si, 0.008 pct P, and 0.006 pct S. The amount of isothermal transformation was determined by electrical resistance measurements using a Kelvin double bridge. The specimens were in the form of rods 1/16 in. diam x 4 in. long. Aus-tenitizing was performed at 1095°C for % hr either in a pre-purified nitrogen atmosphere or in an evacuated Vycor tube, followed by oil quenching to room temperature. After this treatment, the samples were completely austenitic, and the transformation studies were conducted at subatmospheric tempera- tures. Lineal analysis and X-ray determinations of retained austenite were used to calibrate the electrical resistance changes in terms of the percentage of martensite. The martensite range curve for this alloy is given in Fig. 9 of ref. 6. The M, temperature occurs at about -18°C (255°K). Reaction-Path Theory The nucleation and growth description of the martensitic transformation that occurs isothermally has already been given by Kurdjumow.&apos;,&apos;2 However, no comparable treatment exists using the reaction-path concept." Consequently this section presents the reaction-path description of the isothermal mode of the martensitic transformation. According to the latter theory, there is a reaction path (a strain) which describes the relative motion of the atoms during their transference from the austenitic to the martensitic state. The possibility exists, therefore, that the nucleation of a martensitic plate is accomplished by an embryo that comprises a state intermediate between that of austenite and martensite. This state may be attained within a finite volume by means of a homogeneous strain of the austenitic matrix. There is a free-energy barrier (F.) along the reaction path, corresponding to a critical strain in a critical volume. When activation is achieved, a spontaneous increase in strain and volume of the embryo takes place to generate the full-size martensitic plate. This mode of nucleation is to be distinguished from the more conventional mode involving embryos of the martensitic phase that merely grow in size by an interface motion. There are two ways whereby embryos in the austenitic phase can achieve the critical state of strain for isothermal nucleation: 1—through the formation of internal strains by plastic deformation or by stress; and 2—through thermal fluctuations, preferably superimposed upon existing internal strains. The activation free energy F,, must be supplied either mechanically by internal strains or thermally by energy fluctuations, or by a combination of both.

Jan 1, 1953

By Harrison Everett Ashley

Slimes are usually defined as all material passing a certain sized sieve, which is invariably the finest sieve employed by each metallurgist in his tests; 100-mesh and 200-mesh have been taken as the limits by different writers in recent volumes of the Transactions; 150-mesh is common practice. The reader of the first two volumes of Richards&apos; is impressed by the inflexibility with which all grinding-arrangements are praised or condemned according as they make little or much slime. Until recently, one of the principal aims of the mill-man was to keep down the slimes. He was obliged to have separate apparatus for sand- and slime-treatment. The latter was sometimes puzzling and erratic. Often it was found most economical to omit treatment of the slimes entirely, or to store them in immense basins or heaps, waiting for possible future improvements in methods of treatment. Recently, improved processes of agitation and filtration have made possible the successful extraction of values from many slimes. The tendency, when the metal is exceedingly finely disseminated, is to cut out all sand-treatment and to slime everything. These processes have been almost entirely mechanical in their conception, and it is the object of the present paper to show that certain chemical considerations may greatly hinder or assist their successful operation. It has long been a source of dissatisfaction to chemists that their analyses fail to be of much help to the clay-worker; that it is a matter of common occurrence for clays of closely identical ultimate chemjcal constitution to show the widest possible divergence in physical properties.

Jan 1, 1911

By J. G. Savins

Certain types of macromolecules added to water and salt solutions flowing in turbulent motion can reduce the pressure gradient. Alternatively, the volumetric capacity of a pipe for these fluids is increased by the presence of these materials. Examples presented show that the drag reduction can become significant. Thus, the presence of 0.28 per cent of a gum derivative in a solution of sodium chloride flowing at 200 gal/min in a 1.89-in. pipe yields a pressure drop which is 0.44 of the single-phase drop measured under the same conditions of turbulent flow; the addition of 0.1 per cent of a vinyl derivative to a I-in. water line yields a throughput capacity which is 1.78 of the single-phase capacity at the same pressure drop. It is further shown that these phenomena are distinctly different from previous observations with other classes of non-Newtonian systems. There a simple lowering of friction factors below the levels predicted from the resistance laws for Newtonian fluids is associated with a suppression of turbulent motion. A rational physical explanation for drag reduction is advanced. Briefly, the proposed mechanism is a storage by the molecular elastic elements of the macromolecules in solution of the kinetic energy of the turbulent motion. INTRODUCTION This study was inspired by a recent review of some paradoxical drag reduction phenomena in turbulent pipe f1ow.I Under very moderate conditions of turbulent flow, the pressure gradient necessary to pump solutions containing certain specific kinds of polymers, fibers and metallic soaps may become appreciably lower than that required* to pump the solvent, i.e., water or a low-viscosity hydrocarbon, under identical flow rates in the same conduit. As shown by our review, this phenomenon of drag reduction in turbulent duct flow was first noted during the second world war, apparently arising in connection with the development of flame warfare weapons. Since that time several papers illustrating this phenomenon have appeared: Toms, Oldroyd,3 Agoston et al., Bundrant and Matthews, Robertson and Mason,6 Ousterhout and Hall,7 Daily and Bugliarello~ Lummus, Anderson, and Fox. 9 That there are practical applications for techniques which increase discharge or decrease the pressure necessary to transport a liquid through a pipeline is illustrated in the patents which have issued which take advantage of this peculiar phenomenon, e.g., Mysels,l0 Dever, Harbour, and Seifert.11 One also finds fragmentary evidence of this effect in the data pertaining to a few of the polymeric solutions studied by Shaver12 and Dodge.13 However, these investigators were concerned with the development of friction factor vs Reynolds number correlations for a variety of non-Newtonian solutions and suspensions, rather than in a study of drag reduction-. A similar kind of drag reduction effect has been observed in gases. Sproull, for example, reports that adding dust to air flowing in turbulent motion through a pipe results in a lowering of the pressure gradient at identical flow rates. There are also military applications for reducing the drag on hydrodynamic vehicles. For example, the possibility of injecting a rheologically complex fluid into the boundary layers of bodies to reduce the skin friction has been investigated by Fabulal5 and Granville.16 Along somewhat different lines are the drag reduction studies of Kramer.17118 He has shown that skin friction can be reduced by covering the surface of a vehicle with a flexible skin. The effect is apparently due to the boundary layer being stabilized by the presence of the skin. Drag reduction by means of coexisting gas and liquid boundary layers, e.g., film boiling and continuous gas injection, has been proposed by Bradfield, Barkdoll, and Byrne,19 Cess and Sparrow,20&apos;21 Sparrow, Jonsson, and Eckert.22 Here the skin friction occurs between a vapor and a surface rather than between a liquid and a surface. There are several references in the literature to friction-factor correlations for non-Newtonian solutions and suspensions: Shaver and Merri11,23 Dodge and Metzner, Clapp,25 and Thomas.26

Jan 1, 1965

The first of the two fall meetings of the Petroleum Division was held at Los Angeles, Oct 1-2 The attendance was the largest at a Coast meeting of the Division The Institute was officially represented by President Lovejoy, Secretary Parsons and George Otis Smith There were four sessions Thursday afternoon and evening, the latter a dinner meeting (both at the University Club), and Friday morning and after- noon, these two in the auditorium of the Edison building Possibly a dozen papers were presented, all of them constituting a valuable contribution to petroleum technology Hallan N Marsh, Chairman of the Division, presided at the dinner meeting at which Dr Smith addressed the members on "The Kettleman Plan m Retrospect," and C M Nickerson talked on the "Buena Vista Hills Cooperative Development Plan "

Jan 1, 1937

By K. P. Mukherjee, K. P. Gupta, J. Bandyopadhyaya

Even though a lot of work has been done in the past to establish phase equilibrium in the Cu-Si system a re cent investigation casts some doubt about the existence and crystal structure of some of the phases that form in the composition range of 15 to 25 at. pct Si in Cu. The present investigation was carried out using high temperature X-ray diffraction technique along with other standard techniques to study the phases in this composition range. The high temperature 6 phase appears to be tetragonal with parameters a,, = 8.815A, c, = 7.903A, and co/ao = 0.896. The reported bcc E phase exists at room temperature and at least up to 780°C and appears to undergo a transformation near 600°C. The phase appears to be cubic but not of the bcc type. The ? phase appears to undergo a transformation, as has been indicated by earlier investigators, and the low temperature form of .? phase is tetragonal with parameters a, = 7.267A, co = 7.8924, and co/ao = 1.086. THE Cu-Si binary system has been investigated by several investigators1" and several intermediate phases,?,e,?&apos; at lower temperatures and ?,ß,0,e, and ? at higher temperatures, were observed between terminal solid solutions of copper and silicon. Even though the existence of the e phase and the transformation in the ? phase were reported in many early works, in a recent study of this system Nowotny and Bittner6 doubted the existence of the e phase and phase at 550°C. Among the high temperature phases, the 6 phase was reported to have a complex cubic structure with parameter a, = 8.805A.7 Nowotny and Bittner, however, suggested that the structure of the 6 phase might be of CsCl type. In order to check these contradictory reports the present study was taken up to investigate the Cu-Si binary system in the composition range of 17 to 25 at. pct Si. EXPERIMENTAL PROCEDURE Weighed amounts of copper (99.99 pct) and silicon (99.9 pct) were induction melted in recrystallized alumina crucibles under argon gas atmosphere. The alloys containing 17, 18, 20, 21, 21.2, 22, and 24 at. pct Si were annealed in evacuated and sealed quartz capsules at 700°C for 3 days and subsequently water quenched. Other than this annealing, the 21.2 at. pct Si and 24 at. pct Si alloys were annealed at 550°C for 10 days, the 17 at. pct Si alloy was annealed at 750°C for 3 days, and the 22 and 24 at. pct Si alloys were annealed at 780°C for 2 days. All annealing temperatures were controlled to within *l°C. Alloys after quenching were subjected to metallographic and X-ray diffraction investigation. A solution containing 5 g FeC13 + 10 cc HCl + 120 cc H2O diluted with six times its volume with water was used as etching reagent. A 114.6 mm diam Debye Scherrer camera was used for obtaining diffraction patterns. The 17, 21.2, and 24 at. pct Si alloys were subjected to high temperature diffractometry using a Tempress Research High temperature attachment and a GEXRD VI diffractometer. For the 6 phase (17 at. pct Si alloy) powder specimen from a 750°C annealed alloy was reheated to 750°C in the high temperature attachment for 1½ hr before taking a diffraction trace. A 550°C annealed and slowly cooled phase (24 at. pct Si) alloy was first reheated to 550°C. a diffraction trace was made after annealing it for 2 hr, and subsequently it was heated to 716OC and kept at this temperature for 2 hr before taking a diffraction trace. For the e phase (21.2 at. pct Si alloy) a 550°C annealed and slowly cooled specimen was heated first to 425°C and annealed at this temperature for 2 hr before taking a diffraction trace. Subsequently, the specimen temperature was raised to 495", 540°, 603", 635", 682", 720°, and 748°C and homogenized at each temperature for 1 hr before taking diffraction traces. The powder specimen temperature was controlled to within +2oC at each temperature and argon gas, purified by passing it at slow rate through a fused CaC12 column, hot (800°C) copper and titanium chips and finally through a P2O5 column, was used to prevent oxidation of the powder. For all X-ray work copper-radiations at 25 kv, 15 ma (for Debye Scherrer technique), and 40 kv, 20 ma (for diffractometer tech-nique) were used. RESULTS AND DISCUSSION At 700°C the alloys containing 17 to 21 at. pct Si showed two phases while the 21.2 at. pct Si alloy was found to be single phase. The X-ray diffraction patterns of the two-phase alloys were consistent with the phase (ßP-Mn type structure) and the phase (21.2 at. pct Si) patterns. The diffraction patterns of the 17 at. pct Si alloy quenched from 750" and 700°C were identical. According to the accepted Cu-Si phase dia-gram4,5,10 the 17 at. pct Si alloy at 750°C should be in the (k + 6) two-phase region and very close to the -phase boundary. The identical patterns possibly resulted from the decomposition of the 6 phase on

Jan 1, 1970

By C. G. Dunn

It is well known that plastic deformation alters many of the properties of a metal and subsequent heat-treatment partially or completely restores these properties.l In the deformed or strained state, the metal is unstable and tends to change toward a condition called a "strain-free state." The transformation occurs through recovery, recrystallization, and grain growth—processes that may take place singly or in combination. The distortion of the lattice of an individual grain of a metal in a state of strain may be rather complex in nature, because plastic deformation produces: (I) dislocations within mosaic blocks; (2) elastic variations of the lattice spacings; and (3) gross alterations throughout the lattice, especially along slip planes, along composition planes between a grain and its mechanical twins, and along boundaries of deformation bands. These gross alterations are of the nature of bent planes or rotated regions of the crystal lattice and are revealed by a spread in the orientation of the grain. Although we cannot describe these strains and their formation accurately because of insuffcient knowledge, we can, nevertheless, use the information as well as possible to obtain a better understanding of recovery processes. In recovery, the lattice of a grain is not made anew as it is in recrystallization, but is improved or mended in such a way that the basic structure remains unaltered: Until recatly observations were .that recovery produced no marked changes in the shapes of spots in Laue diffraction patterns, whereas recrystallization did, but now it is known for silicon ferrite2 that Laue spots may become quite sharp entirely through recovery. Consequently, the shape of Laue spots alone would not be a suitable test to distinguish between recovery and recrystallization. There is considerable evidence that the micro-structure usually does not change visibly during recovery. Absence of a visible change in the microstructure, therefore, provides a sufficient test of recovery in many cases. However, including this observation in a definition of recovery as a necessary condition (this is usually done) is unfortunate, because recovery may, as will become evident later, produce new grain boundaries that are visible not only in the microstructure but also in the macrostructure. For the present, therefore, let us say that a necessary condition for a process to be one of recovery is that the principal orientation or orientations of a deformed grain be essentially unchanged throughout the transformation toward the strain-free state. Several transformations may occur that fulfill this condition, and the nature of the distortion in a grain indicates what these must be if the lattice is to be mended in part or fully. Consequently, it will be convenient as well as

Jan 1, 1947

By C. G. Dunn

It is well known that plastic deformation alters many of the properties of a metal and subsequent heat-treatment partially or completely restores these properties.l In the deformed or strained state, the metal is unstable and tends to change toward a condition called a "strain-free state." The transformation occurs through recovery, recrystallization, and grain growth—processes that may take place singly or in combination. The distortion of the lattice of an individual grain of a metal in a state of strain may be rather complex in nature, because plastic deformation produces: (I) dislocations within mosaic blocks; (2) elastic variations of the lattice spacings; and (3) gross alterations throughout the lattice, especially along slip planes, along composition planes between a grain and its mechanical twins, and along boundaries of deformation bands. These gross alterations are of the nature of bent planes or rotated regions of the crystal lattice and are revealed by a spread in the orientation of the grain. Although we cannot describe these strains and their formation accurately because of insuffcient knowledge, we can, nevertheless, use the information as well as possible to obtain a better understanding of recovery processes. In recovery, the lattice of a grain is not made anew as it is in recrystallization, but is improved or mended in such a way that the basic structure remains unaltered: Until recatly observations were .that recovery produced no marked changes in the shapes of spots in Laue diffraction patterns, whereas recrystallization did, but now it is known for silicon ferrite2 that Laue spots may become quite sharp entirely through recovery. Consequently, the shape of Laue spots alone would not be a suitable test to distinguish between recovery and recrystallization. There is considerable evidence that the micro-structure usually does not change visibly during recovery. Absence of a visible change in the microstructure, therefore, provides a sufficient test of recovery in many cases. However, including this observation in a definition of recovery as a necessary condition (this is usually done) is unfortunate, because recovery may, as will become evident later, produce new grain boundaries that are visible not only in the microstructure but also in the macrostructure. For the present, therefore, let us say that a necessary condition for a process to be one of recovery is that the principal orientation or orientations of a deformed grain be essentially unchanged throughout the transformation toward the strain-free state. Several transformations may occur that fulfill this condition, and the nature of the distortion in a grain indicates what these must be if the lattice is to be mended in part or fully. Consequently, it will be convenient as well as

Jan 1, 1947

By Leonard Obert

The structural stability of any mine or under- ground opening in rock is dependent on the stress field, that is, the state of stress in rock before mining, the stress distribution in the rock created by the opening, and the in situ strength of the rock. In the past mine design problems have been considered by assuming that the stress field is due only to the weight of the overlying rock, the stress around the opening can be calculated from elastic theory using laboratory-determined elastic constants, and the in situ strength of the rock can be determined from laboratory tests on small specimens. Solutions to hypothetical problems using this procedure are given in a report by Obert, Duvall, and Merrill.&apos; Because of the uncertainties of these assumptions, this method can give only a rough approximation to the solution of actual structural problems. Obviously, this approach to determining the stability of underground openings would be greatly improved if the rock stress could be measured in situ, and it is undoubtedly for this reason that in recent years many investigators have spent considerable effort on the development of procedures for measuring rock stress. In this report a method is described for determining in situ the absolute stress in rock. The method requires the measurement of the deformation of a borehole before and after the rock is stress-relieved. The magnitude and direction of the applied stresses are calculated from the borehole deformation and modulus of elasticity of the rock. Also, the results of absolute stress determination made in the walls and pillars of underground openings in five rock types are presented. From these results a number of conclusions are drawn regarding the accuracy and limitations of this method and the problem of stress measurements in general.

Jan 8, 1962

By Cyril Wells, Robert F. Mehl

THE practical importance of a knowledge of the rates of diffusion of carbon and of alloying elements in steel has often been pointed out 1,13 This importance ties chiefly, though by no means only, in the study of carburizing and in the study of reactions involved in the eutectoid inversion in steel The present paper is one of a series in which the rates of diffusion of carbon and of alloying elements in y-Fe and their mutual effects are measured with precision The present paper reports the results of meas¬urements of diffusion coefficients D of manganese in -y-Fe over wide ranges of temperature, and concentration, and of the effect of carbon and impurities and of grain size upon these diffusion coefficients [$] Data on the rate of diffusion of manganese in y-Fe have been reported by Fry2 and by Paschke and Hauttmann 3 Claussen4 calculated the diffusion coefficient D from Fry&apos;s data as 3 X 10-10 sq cm per sec at 960°, which will be shown to be incorrect by many orders of magnitude Paschke and Hauttmann gave a rough value of 9 6 X 10-8 sq cm per sec at 1400° Neither of these workers considered the variation of D with concentration and with temperature, and neither studied the effect of carbon METHOD AND MATERIALS Method -The method employed is identical with that used earlier 5 Owing to the volatility of manganese and the consequent loss of manganese from the outer surfaces of the welded diffusion bars, these surfaces were discarded The diffusion anneals were performed in argon, except one (test 18) in which hydrogen was employed for purposes noted below The temperatures reported are accurate to about 3° and constant within I°C , the diffusion times are reported to the nearest half hour The specimens were furnace-cooled after the diffusion anneal at about 2° per min Layers, usually 0 004 in thick, parallel to the welded interface, were machined from the specimens after the diffusion anneal and analyzed for manganese to 0 01 per cent Mn, the distances of these layers from the interface was measured to ±0 0002 in The concentration-penetration curves obtained were analyzed by both the Grube and the Matano solutions e The D values calculated by these two methods do not differ appreciably 5 The Grube solution was used except in tests in which the effect of concentration upon D was greater than 50 per cent, in such cases the Matano method, which explicitly provides for the variation of D with concentration,1 is the more reliable In general, the number of layers machined from the zones of variable composition were large enough to permit the use of the average layer composition as that of the center of the layer In 3 of the 30 tests, however, this is not justified, and the D values

Jan 1, 1941

SPEAKING generally, it may be said that leaching is the simplest method of recovering copper from its ores. Likewise it is perhaps the oldest method of treatment used by copper metallurgists of the modern era. Reliable records indicate that as early as 1752 cupriferous pyrite was leached in the famous Rio Tinto region in Spain. Though the manipulation was quite different from that in use today at Chuquicamata, New Cornelia, Inspiration, and Andes, it resembled that employed by Phelps Dodge at Bisbee, as described in Chapter XIV. Leaching consists of subjecting ore to the action of a liquid solvent or menstruum (usually an aqueous solution) which dissolves the major part of the soluble copper minerals. The resulting pregnant solution is drained off or otherwise separated from the residual ore; and the copper in the solution is then recovered by some appropriate means of precipitation. When sulphuric acid is the lixiviant, as is usual, the simplest procedure is to pass the copper-bearing solution through large troughs containing scrap iron. The iron displaces the copper in the solution, precipitating it as a slime coating the iron and varying in color -from bright copper to a dull black. This may be washed off at intervals, collected as a sludge, dried, and melted with a little fluxing material to produce reasonably pure copper. Usually, of course, the mud is sent to a smelter where most often its destination is the converter or reverberatory furnace. Modern practice uses precipitation on iron only as a complemental process to electrolytic precipitation or deposition, wherein the pregnant copper-bearing solution, acting as the electrolyte in a cell having insoluble anodes and copper cathodes, is depleted of the copper it contains. Cathode copper, 99.5 per cent pure, is produced ready for the market as refined copper, or for a simple process of melting, fire-refining, and casting in

Jan 1, 1933