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Both drilling and production of oil decreased in Rumania during 1935. The drilling totaled 1,042,041 ft. and production 61,973,938 bhl., being a decrease of 8.5 and 1.6 per cent, respectively. Wells completed during the year totaled 198, of which 168 were oil wells, compared with 238 and 215, respectively, of 1934. No important discovery was made during the year. The only new field found was Valea Calugareasca, where the Credit Minier completed a small well on the westward plunge of the anticline. Drilling higher on the structure will probably prove the field to be of greater importance than the results of the discovery well indicate. The only other developments of importance were the extension of the Bucsani field, the opening up of a new area on the northeast of the salt at Moreni by Credit Minier and a minor extension of the West Gura Ocnitei field. At Calinesti, on the cast extension of the Moreni field, exploration work continued and a gas well was completed, but oil in commercial quantities is yet to be found.

Jan 1, 1936

Both drilling and production of oil decreased in Rumania during 1935. The drilling totaled 1,042,041 ft. and production 61,973,938 bhl., being a decrease of 8.5 and 1.6 per cent, respectively. Wells completed during the year totaled 198, of which 168 were oil wells, compared with 238 and 215, respectively, of 1934. No important discovery was made during the year. The only new field found was Valea Calugareasca, where the Credit Minier completed a small well on the westward plunge of the anticline. Drilling higher on the structure will probably prove the field to be of greater importance than the results of the discovery well indicate. The only other developments of importance were the extension of the Bucsani field, the opening up of a new area on the northeast of the salt at Moreni by Credit Minier and a minor extension of the West Gura Ocnitei field. At Calinesti, on the cast extension of the Moreni field, exploration work continued and a gas well was completed, but oil in commercial quantities is yet to be found.

Jan 1, 1936

The following rules, upon approval by the Board of Directors, shall serve as guides in the practical applications of the Bylaws for the operation of the Institute and have the full power and effect of Bylaws, however, these rules may be changed, modified, re¬pealed or amended when deemed necessary by the Board of Directors Any change or alteration of these Rules shall be published in the Journals of the Institute within three months after such alteration is adopted ARTICLE I-ORGANIZATIONS OF THE INSTITUTE Section 1-Societies Rule 1-Governing Body Each Society shall form its own governing body composed of an Executive Committee, or similarly designated executive body, duly elected officers, and a Nominating Committee for the nomination of Institute officers as provided by the Bylaws of AIME and of each Society approved by the Board In all correspondence and printed matter the Society Boards of Directors shall be specifically identified For ex¬ample, the Mining Board of Directors or the Board of Directors of the Society of Mining Engineers Using simply the "Board of Directors" could cause confusion

Jan 1, 1960

By Francis T. Havard

A Reply to the Discussion by Mr. Francis T. Havard of the Paper by Mr. Henry W. Edwards, which was read at the Atlantic City Meeting, February, 1.904. (Annual Meeting, February, 1905.) MR. HENRY W. EDWARDS, Grand Junction, Colo. (communication to the Secretary) :* Effect of Heat. It should be remembered that the coefficient of expansion of concrete is nearly that of iron; therefore in constructing flues the design must take account of this important detail. The cracking of flues, mentioned by Mr. Havard as having occurred at a temperature of 100° C. and above, . were possibly caused by neglecting to appreciate the extent of the expansion and contraction of the concrete. Effect of Atmospheric Moisture.-The moisture present in the atmosphere is exceedingly injurious to flues of any material, and those made of concrete are no more susceptible than of brick. My experience with flues of the Monier construction, having walls from 1.5 to 2.5 in. thick plastered upon a net-work of iron rods, was unsatisfactory from the standpoint of dura¬bility.

Mar 1, 1905

Geological Survey of Alabama, University, Ala. A complete list of publications will be sent upon request Requests for publications should be addressed to Walter B. Jones, State Geologist, University, Alabama Publications of the Survey, of most interest to the general public, comprise (1) Special Reports, (2) Monographs, (3) Bulletins, (4) Circulars, (5) Maps The Special Reports mostly deal with a special coal field, mountain, valley, or series of formations, while the Monographs treat of some industry, resource, or group of organisms throughout the state or as far as it extends The Bulletins deal with a variety of subjects, but are usually smaller and less technical than the Monographs and Special Reports Maps now available comprise a map of the Cahaba coal field, one of the Coosa coal field and a small geological map of the state, in addition to a large geological map, 28 by 42 inches, 8 miles to the inch (1926) (also bound with Special Report 14) Price of the large geological map of Alabama, folded in manila envelope, is $1.03 postpaid

Jan 1, 1933

By E. E. ELLIS

THE zinc deposits under development by the Universal Exploration Co., a subsidiary of the United States Steel Corporation, are close to Jefferson City, Tenn., a small town about 30 miles northeast of Knoxville, Tenn., and are served by the Southern Railway. The first zinc mining in this district was in the old Mossy Creek mine at the north edge of the town, opened for oxidized ores in 1871 and later worked for sulfide ores. An oxidized zinc ore deposit (carbonates and silicates) was opened near New Market in 1900 and worked through the war period by the G1'8sselli Chemical Co., producing around 25,000 tons of oxidized ore. In recent years a sulfide deposit has been found and developed but is not producing. About 1% miles south of this a small body of sulfide ore was developed by open cut, but the operation was unsuccessful. The Universal Exploration Co. entered the district in 1926 and, after extensive exploration, has been operating an oxidized ore-body since Janual:y, 1928, and has been developing a sulfide ore body since July, 1928.

Jan 1, 1929

By George E. Beavers

It is well to understand at once that converter gas is always mixed with copper blast-furnace gas at Copperhill before entering the Glover tower, and, at times, roaster gas is also added along with the converter and blast-furnace gas. Fig. 1 is introduced to show the variations in the sulfur dioxide content of the gas and the ability of the acid-plant operator to control the variations, resulting in a very even grade at the Gay-Lussac tower entrance. The sulfur dioxide content of the gas entering the Gay-Lussac tower must be low and regular to assure low niter consumption. The blast-furnace gas alone is rather uniform as regards sulfur dioxide content and the variations, as shown in Fig. 1, are chiefly caused by converter changes—out for charging, slag pouring, copper pouring, and, of course, the "finish." The combined gas during the converter finish is, at times, almost free of oxygen and the sulfur dioxide content reaches 13.0 per cent. At such times air is added at the Glover tower.

Jan 1, 1934

By V. C. Allen

THE Tri-State Zinc, Inc., Galena, Ill., was confronted with the problem of securing ore from a deposit because the hoisting shaft was several thousand feet from the mill. The orebody is several thousand feet long, averaging 200 ft in width and 60 ft in height and opened up by vertical shafts some 300 ft deep. Mining is by the room-and-pillar method. During the initial operation the ore was loaded by conventional electric 1/2 -yd boom-and-dipper shovels and hauled to the shaft by 8-ton diesel trucks. This underground ore loading and hauling was well adapted to the conditions and productive of low costs per ton. However, with the mill situated as mentioned, a triple handling of all broken rock was necessary: 1-from the stope to the shaft by truck, 2-up the shaft by skip or can into the surface hopper, and 3-by truck from the surface hopper to the crushing plant at the mill. In addition to the repeated handling, serious troubles were encountered during the winter because of freezing in the shaft hopper.

Jan 6, 1951

By N. J. Grant, F. B. Cuff

DURING a recent investigation of high-purity aluminum,' evidence was found for grain-boundary sliding and migration at room temperature. The investigation was concerned primarily with an electron-microscopic study of a fine (approximately 1 micron) network that appeared on the surface of chemically polished aluminum. The nature of the network will not be discussed since it serves only as markers showing relative displacements in this instance. The pertinent observation, shown in Fig. 1, was made on a flat tensile specimen which had been strained about 8 pct at room temperature. As can be seen in this electron micrograph, a grain boundary has undergone a certain amount of sliding. The reason for the existence of the two differently appearing networks is given in the foregoing reference. For present purposes, however, the transition region can be disregarded. A close study of this picture reveals that in addi- tion to sliding there is a narrow band adjacent to and above the grain boundary. Fig. 2 is a schematic drawing showing the essential features of this band.

Jan 1, 1959

By M. J. Marcinkowski, K. S. Raghavan, A. S. Sastri

A detailed transmission electron microscopy investigation was carried out in connection with a manganese Hadfield Steel. At small plastic strains, numerous individual intrinsic stacking faults are observed. With increased plastic deformation, the stacking faults thicken into twin lamellae which in turn subdivide the original austenite matrix into smaller domains. The twin boundaries act as strong barriers to subsequent dislocation motion and is in a sense equivalent to grain refinement. It is this "grain refinement" which is believed to be the cause of the very high work hardening rates in the Hadfield Steels. In many cases, especially where an hcp phase is the stable one at low temperatures, the stacking fault energy in fcc metals and alloys decreases with decreasing temperature.' Since stacking faults of the intrinsic type are precursors of both twins as well as the hexagonal close packed structure, both of these entities should become more frequent as the temperature of a fcc crystal is lowered. In the case of the twin, there is no chemical driving force for its formation and it is generally necessary to provide the required driving force by an applied stress, i.e., strain energy. In the case of the hcp structure the transformation from the fcc modification can occur spontaneously (marten-sitically) since a decrease in chemical energy does in fact occur; however, an applied stress will provide an even greater driving force toward complete transformation. Since the transformation products mentioned above occur in an inhomogeneous manner throughout the crystal and since these can act as potential barriers to further plastic deformation2 marked strengthening effects can be anticipated. Also because metal and alloy strenghening is in general proportional to the shear modulus, these effects should be greatest in steels of the austenitic type (y), i.e., the fcc types. Perhaps the two most important steels in this category are the austenitic stainless steels and the Hadfield manganese steels. Both may be quenched from elevated temperatures so as to retain the austenitic states characteristic of those temperatures. The effect of subsequent deformation at lower tem- peratures has a profound effect on the stress-strain curves of these alloys. In particular Fig. 1 shows the compressive stress-strain curves obtained with an 18-8 stainless steel which was quenched from 1850°C after annealing for 1 hr so as to produce all y. As the temperature is lowered, the work hardening rate increases markedly. Although some hcp or c mar-tensite can be generated by plastic deformation as the temperature is lowered,~ it is believed to be a transition phase4 and most of the martensite produced is of the bcc or a variety.3 It is this stress induced martensite which gives rise to the very low initial work hardening at 77°K as can be seen in the stress-strain curve in Fig. 1. Similar low initial work hardening rates have been observed in the stress induced Ni-Ti martensites.5 Fig. 2 shows that an even more rapid rate of work hardening occurs in the Hadfield steels treated in the same way as that described for the 18-8 stainless steels a; the temperature is lowered. It is this ability to work harden to such high stress levels that makes the Hadfield steels particularly suitable for armor plate and heavy construction equipment. However, unlike the case of Fig. 1, no initial low rate of work hardening is observed in any of the curves in Fig. 2. Thus the stress induced formation of any low energy martensite phase in any significant quantity must be ruled out. This observation is in accord with the X-ray findings of Otte.~ On the other hand, small quantities of the E phase have been observed by other investigators using transmission electron microscopy (TEM) above Even more significant was the fact that large numbers of deformation twins were observed in the deformed Hadfield steels,678 which were postulated to be one of the reasons for the high work hardening ability of this class of steels.8 It is the purpose in what follows to discuss a series of experimental observations pertaining to the stress induced transformation in a Hadfield steel and to formulate a dislocation mechanism which adequately accounts for the observed results. EXPERIMENTAL PROCEDURE The stainless steel used to obtain the curves shown in Fig. 1 was of the AISI Type 303 containing approximately 18.0 pct Cr and 8 pct Ni. On the other hand, the Hadfield manganese steel used to obtain the curves shown in Fig. 2 contained between 1.00 to 1.25 pct and 11.5 to 13.5 pct Mn. In all cases the samples were in the form of compression cylinders 0.220 in. in diam and 0.370 in. long. Prior to testing the samples were annealed for a hr at 1050°C and rapidly quenched into a brine solution. This treatment was sufficient to preserve the y phase for subsequent testing at lower temperatures. All samples were compressed in an Instron testing machine using a cross head speed of 0.02 in.

Jan 1, 1970

By Harold M. Feder, Robert V. Schablaske, Irving Johnson, Ewald Veleckis

The stoichiometry, structure, and stability of the internzediate phases formed between zinc and some of the rare earth (RE) metals were systematically exarnined by means of a recording effusion balance and X-ray diffraction analyses. In the La-, Ce-, PY-, Nd-, and Y-Zu systems, at or below about 600 C, the following sequences of phases (REZnx) were found: La, x = 1, 2, 4.0, 5.25, 7.3, 17/2, 11, and 13.0;' Ce, x = 1, 2, 3, 11/3, 4.3, 5.25, 7.0, 17/2,* and 11; Pr,x = 1,2, 3, 11/3 ,* 4.3, 5.3(?), 7.0, 17/2,* and 11; Nd,x = 1,2, 3,* 11/3,* 4.3, 6.5, 8.5,* and 11; Y,x = 1,2,3, 11/3, 4.5, 5.0, 17/2,* and 12.* The structure types of all these phases were classified. In addition, lattice parameters were obtained for the first time for the pluses denoted by asterisks. In the absence of de tectable valency or electronegativity effects the systesnatic trends in the results have been ascribed to the effects of' the lanthanide contraction. For example, the maximum number of zinc atoms in the coordination polyhedron surrounding the RE atom decreases from twenty-four to twenty-two to twenty as the size of the RE atom decreases. THE structures and compositions of a great many intermetallic phases (e.g., the Laves phases) are known to be based primarily, but not exclusively, on the space-filling efficiency of various modes of packing together atoms of different sizes. The valencies and electronegativities of the constituent atoms are, however, also influential. In extreme cases hypothetical intermetallic phases which fulfill the efficient spacefilling requirements may not be present in the constitutional diagram because of thermodynamic instability brought about by the operation of valency or electronegativity factors. Hence, for a detailed study of the influence of atomic size on alloy structure and composition, it would be desirable to minimize variations of valency and electronegativity. The intermetallic phases formed by the rave earths (RE) with some common partner offer an excellent opportunity for isolating the effects of size from those of valency and electronegativity. The rare earths exhibit a large, but smooth, decrease in size (the lanthanide contraction) in the series from lanthanum to lutetium when inter comparison is made for a common valence state, e.g., isolated atoms or trivalent ions. The elements yttrium and scandium are frequently included as pseudo rare earths; their sizes place them in the vicinity of dysprosium and lutetium, respectively. The electronegativities of RE elements vary by less than 10 pet. The trivalent state is the most common; however, the well-known tendency of cerium, praseodymium, and terbium to achieve higher valencies, and of samarium, europium, and ytterbium to seek lower valencies, requires that caution be exercised in the assumption of equal valencies. In the present study the existence, constitution, and structure of each of the numerous intermediate phases formed by zinc with lanthanum, cerium, praseodymium, neodymium, or yttrium were examined systematically and in detail. The investigation was conducted by a recording effusion balance technique and by X-ray diffraction analysis. The results enrich our knowledge of the phase diagrams of these systems. In addition, they present some clear-cut evidences of the operation of the size factor alone. EXPERIMENTAL PROCEDURE Apparatus. The mode of operation of the recording effusion balance and its application to phase studies have been discussed in detail elsewhere.' In this work, an effusion cell containing a finely divided alloy was suspended within an evacuated tube from the beam of an analytical balance. The tube was immersed in a massive molten salt bath whose temperature was controlled to within 0.5o C during each experiment. The loss in weight of the alloy owing to effusion of zinc* was continuously recorded. Two effusion cells, 1/2 in. diam by 1 in. high, were machined from tantalum rods. Two orifices were drilled laterally into the walls of each cell. The orifice areas were determined by calibration with pure zinc: cell A had a total orifice area of 6.5 x 10-41 sq cm, and cell B an orifice area of 9.8 x 10-3 sq cm. By appropriate choices of orifice area and temperature the wide range of volatilities from pure zinc to pure rare earth metal could be investigated. X-ray diffraction powder photographs were made at room temperature with a 114.6-mm Debye-Scherrer camera with both filtered CuKa radiation and filtered CrKa radiation. Lattice parameters were refined by a computer-programmed least-squares analytical treatment which incorporated appropriate extrapolation techniques.2 Frequent use was also made of a special computer program3 designed to generate a powder pattern from an assumed structure in order to verify structural assignments. Materials. Lanthanum, neodymium, and yttrium were purchased from the Lunex Co., cerium from the Cerium Metals Corp., and praseodymium from the St.

Jan 1, 1968

By H. Conrad, V. V. Damiano, G. J. London

A method is described for producing single crystals of high-purity beryllium, Be-4.37pct Cu, and Be-5.24 pct Ni. These crystals were prepared for testing in compression parallel to the [0001] by orienting and lapping to within ±3' of arc of the (0001). Microstrain testing apparatus is described along with c axis compression results for ingot purity beryllium, twelve-zone-pass material, and the above-mentioned alloys. Results show no measurable plasticity for the ingot purity material from -196" to 400°C, although some surface traces of (1122) slip was observed at 200°C and above. The twelve-zone-pass material shows substantial microstrain plasticity at 220°C with slip on (1122). Both alloys show significant plasticity at room temperature and above with slip also on (1122) planes. THE two slip systems which normally operate during the plastic deformation of beryllium in the vicinity of room temperature are:' basal slip (0001)(1120) and prism slip . Pyramidal slip with a vector inclined to the basal plane has been reported for elevated temperatures,'-a but occurs near room temperature only at very high stresses.~ A summary of the available data on the effect of temperature on the critical resolved shear stress for slip on these systems has been compiled by Conrad and Perlmutter.~ It has been postulated6'7 that one of the principal factors contributing to the brittleness of poly crystalline beryllium at temperatures below about 200°C is the difficulty of operating pyramidal slip with a vector inclined to the basal plane. Hence, detailed information on the operation of such a slip system is important to understanding the brittleness of beryllium. The operation of pyramidal slip with a vector inclined to the basal plane is best accomplished in beryllium by compressing single crystals in a direction parallel to the c axis. In such a test the resolved macroscopic shear strzss on the basal and prism planes is zero and (1012) twinning which is favored by tension along the c axis does not occur. Hence, in c axis compression of beryllium the normal deformation modes are inhibited and the operation of pyramidal slip with a vector inclined to the basal plane is favored. In the present investigation, c axis compression tests were performed on beryllium single crystal as a function of temperature (77" to 700°K), purity (commercial and twelve zone pass), and alloy content (4.37 wt pct Cu and 5.24 wt pct Ni). Presented here is a description of the test techniques employed and the gross mechanical behavior observed. A detailed analysis of the slip traces developed on the surfaces of the deformed specimens during these tests and the results of electron transmission studies of the deformed crystals are given in a separate paper.B PROCEDURE 1) Materials and Preparation. Single crystals about 1 in. diam were prepared of the following materials: commercial-purity beryllium, high-purity beryllium, and two beryllium alloys, one with 4.37 wt pct Cu and the other with 5.24 wt pct Ni. The commercial-purity single crystals were obtained by cutting specimens from large-grained ingot of Pechiney SR material, which is approximately 99.98 pct pure. The high-purity crystals were prepared by floating-zone refining (twelve passes) a rod (7 in. by 1 in, diam) of Pechiney SR grade cast and extruded beryllium. Although an absolute chemical analysis of the zone-refined material was not established, mass spectro-graphic analysis, emission spectrographic analysis, and y activation analysis indicated that it contained in atomic fractions about 5 to 10 ppm each of carbon and oxygen, 1 to 5 ppm each of nickel and iron, and about 1 to 2 ppm of copper, with the remaining residual impurities being less than 1 ppm. Further indication of the purity of this material is provided by the critical resolved shear stress for basal slip, which was approximately 300 psi. The starting material for the alloy single crystals was 1-in.-diam floating-zone-refined (six passes) rod of Pechiney SR grade beryllium. Two such rods were wrapped respectively with sufficient weight of wire of high-purity copper (99.999 pct) or nickel (99.999 pct) to yield a 5 wt pct alloy. A seventh floating-zone pass was then applied to each of the rods to accomplish the initial alloying and an eighth pass for homogenization. Analytical samples were taken from regions of the rod immediately adjacent to where the mechanical test specimens were cut; these indicated 4.37 wt pct Cu and 5.24 wt pct Ni. 2) Crystal Orientation. To avoid the occurrence of basal slip during c axis compression testing, it is necessary to load the crystals as nearly parallel to the c axis as possible. Preliminary c axis compression tests indicated that plastic flow and/or fracture occurred at stresses of the order of 300,000 psi; hence on the basis of a critical resolved shear stress for basal slip of 300 to 400 psi, the maximum crystal misorientation permitted is about 4 to 5' of arc. Since this accuracy cannot be obtained using the usual back-

Jan 1, 1969

By Robert D. Pehlke, Robert G. Blossey

The solubility of nitrogen in liquid binary and ternary Fe-Ni-Co alloys has been measured by the Sieverts' method between 1550°and 1700°C. Solubility data and standard free energzes and enthalpies of solution for nitrogen in the alloys are presented. Interaction parameters are discussed and presented for binary and ternary alloys. MOST of the studies of nitrogen solubility in liquid metals have been directed toward the dilute alloys of iron. Several of these investigations have included measurements of the nitrogen solubility in Fe-Ni al10s'- and in Fe-Co alloys.435 There has been some work, however, that has extended across the e-i-" and F-CO" binaries. This investigation was undertaken to determine the nitrogen solubility in both binary and ternary alloys of the Fe-Ni-Co system. It was also hoped that the differences between earlier studies might be resolved. EXPERIMENTAL METHOD This investigation was made using a Sieverts' apparatus described previously." The nickel (99.85 pct) and cobalt (99.9 pct) were obtained from Sherritt-Gordon Mines, Ltd., and the iron (99.95 pct) was Fer-rovac-E obtained from Crucible Steel Co. Recrystal-lized alumina crucibles were used throughout the entire investigation with no evidence of crucible-melt reaction. Melt temperatures were measured with an optical pyrometer and the temperature scale calibrated against the melting points of the three pure metals. The emissivity of the melt was assumed to be a linear function of composition for all alloys, as has been shown for Fe-Ni alloys.lZ The emissivity of the pure metals at 1600°C were taken as 0.43 for iron, 0.44 for cobalt, and 0.45 for nickel. Using these emissivities, the trans mis sivity of the system was found to be 0.51 i 0.01. The Sieverts' method was used for this study and followed the procedures outlined previously.l' The individual metals were weighed to give about 100 g of alloy. The alloys were melted in the crucible under a partial pressure of argon. The system was evacuated, and the "hot volume" was measured with argon. To avoid the errors caused by vaporization, the melt was held under vacuum only long enough to ensure that all of the gas in the system had been removed. The influence of any small amount of vaporization on the "hot volume" was shown to be negligible by measuring the "hot volume" after a run. This measurement agreed with that made at the start of the run within the implicit error, 0.2 cc, caused by the limitations in accurately reading the buret. The solubility-pressure relationship was measured in the pure metals and in several alloy compositions throughout the ternary system. These measurements were made by admitting measured amounts of nitrogen to the system, and then determining the equilibrium nitrogen pressure above the melt. This method has the distinct advantage of higher accuracy, particularly at lower pressures, than measurements made by withdrawing gas from the system to reduce the pressure after determining the solubility at 1 atm nitrogen pressure. This latter method has a practical lower limit of about 0.4 atm where an increased error is encountered because the buret must be emptied to permit further measurements at lower pressures. By determining the relation between apparent solubility and pressure, it was possible to make a good estimate of the initial nitrogen content of the metal from the intercept of the solubility curve extrapolated to zero pressure.11 DISCUSSION The solubility data corrected to 1 atm nitrogen pressure are summarized in Table I. The reported solubility has been corrected for the initial nitrogen content of the alloys. The initial nitrogen contents fell between 0.0002 and 0.0010 wt pct, and were lower in the iron and nickel than in the cobalt. Sieverts' law was obeyed in all alloys at pressures up to 1 atm. Examples of this behavior are shown in Fig. 1. The reaction for solution of nitrogen is Taking the standard state as 1 wt pct N in the alloy and the reference state as nitrogen at infinite dilution in the alloy, and noting the adherence to Sieverts' law, K becomes the solubility of nitrogen in the alloy at 1 atm pressure. Thus the solubility data of Table I were used directly to calculate the standard free energy for the solution reaction. These results are also presented in Table I. The enthalpy of solution is also summarized in Table I as calculated from a form of the van't Hoff relation: Iron-Nickel System. The data for the solubility of nitrogen in liquid Fe-Ni binary alloys is presented in Fig. 2 along the with data of aito, Schenck et al.,' and Humbert and 1liott.l' The data for studies of nitrogen solubility in Fe-Ni alloys containing less than 20 pc t i'- are not presented in Fig. 2, although they are in good agreement with the present work. The results of this study are in good agreement with Schenck

Jan 1, 1967

By C. S. Roberts

The creep mechanism and kinetics of fine-grained magnesium have been studied over the temperature range 200&apos; to 600°F. As a result of a photographic study of microstructural changes, transient and steady-state creep components have been correlated with slip, subgrain formation, and cyclic deformation at the grain boundaries. THE approach of this research has been the blend of a quantitative study of the creep strain of polycrystalline magnesium as a function of time, stress, and temperature with direct microstructural observations of the operative deformation processes. The validity of the conclusions is dependent on the condition that the microstructural changes seen on the polished surface qualitatively represent those occurring in the bulk of the metal. The work was intended as much to lay a background to a study of highly creep-resistant magnesium alloys as to provide a description of the behavior of the base metal itself. The spectroscopic analysis of the electrolytic magnesium used in this study is as follows: Al, 0.009 pct; Ca, <0.01; Cu, 0.0011; Fe, 0.021; Mn, 0.012; Ni, 0.0004; Pb, 0.0012; Si, <0.001; Sn, <0.001; and Zn, <0.01. The impurity level is approximately that of commercial magnesium alloys. The original ingot was melted under Dow type 310 flux and cast as a 3 in. diam billet. It was extruded into 1 in. flat stock under the conditions: billet preheat 800°F (1 hr), container and die temperature 800°F, speed 3 ft per min, and area reduction ratio 45:1. The extrusion process was chosen in preference to rolling and recrystallization because it allowed easier grain size control from specimen to specimen. The grains of the extruded metal were fairly equi-axial and uniform in the size range of 4 to 6 thousandths of an inch. The preferred orientation of basal planes about the transverse direction was determined by an X-ray diffraction surface reflection method. A beam of filtered copper radiation was directed at an angle of 17" to both the transverse direction and the surface yet perpendicular to the extrusion axis. Analysis of the (002) diffraction arcs in the resulting photographic patterns gave an approximate intensity distribution along the great circle which extends through the center of the basal plane pole figure and to the extrusion axis poles. Successive layers of metal were removed by macro-etching between exposures. The extruded texture is relatively sharp, but the most significant point is the position of the maximum basal plane pole density and its variation with depth below the surface. Fig. 1 shows that this maximum is rotated 15" from the normal at the surface toward the extrusion direction. Such an inclination has been reported for extruded 1 pct Mn and 8 pct A1-0.5 pct Zn alloys.&apos; The inclination decreases until the maximum splits at about 0.025 in. depth into two elements of equal and opposite rotations from the ideal. The double texture persists to as great a depth as was experimentally convenient to examine. It probably continues to the very center of the extrusion. There is no great change in the sharpness of the individual elements of the texture with depth. A plate of metal about 0.015 in. thick at the surface of the extruded stock was produced by etching. A transmission diffraction pattern was made for the purpose of determining any preferred orientation of a direction in the basal planes. Relatively uniform {loo) and {101) rings were produced. There is little tendency for parallelism of a given direction in the plane with the projection of the extrusion axis on it. The creep specimens were machined from 6¼ in. lengths of the extruded stock. Creep was measured on the reduced section, ½x1/8X2¼ in. long. This section was electropolished on one side for the studies of microstructural changes during creep. An orthophosphoric acid-ethyl alcohol electrolyte was used under the conditions recommended by Jacquet.² Hand polishing was used for previous mechanical preparation. Electropolishing was continued until all mechanical twins had been removed. The electro-polished surface was protected from oxidation during creep testing by a thin layer of silicone oil. All micrographs were taken at room temperature on conventional metallographic equipment and after removal of the oil film. The creep tests were performed with machines which have been described in detail by Moore and McDonald." Five testing temperatures, 200°, 300°, 400°, 500°, and 600° ±3°F were used. Difference in temperature between the two ends of the specimen reduced section was 2°F or less. The testing was done at constant load. Strain readings were taken as frequently as necessary to develop usable creep curves. Tensile Creep vs Time, Stress, and Temperature A definition of terms is necessary. Whenever successive sections of a creep strain-time curve show decreasing, constant, and increasing slope with time they will be termed primary, secondary, and tertiary

Jan 1, 1954

By Pradeep K. Rohatgi, Clyde M. Adams

Structures produced upon solidification of the eu-tectic composition (33 wt pct Cu) aluminum copper alloy have been examined as a function of freezing rate dfs /d? , the rate of change of fraction solid (fs) with time (8). Slow (dfs/d? = 0.0016 sec-1), intermediate (dfs/d? = 0.02 sec-1) and rapid (dfs/d? = 0.4 to 7.30 sec-1) freezing rates were used. The lamellar Al-Cual2 eutectic is arranged in the form of rod-shaped colonies at rapid freezing rates. The colonies are aligned parallel to the direction of heat flow, whereas the lamellae within the colonies are aligned at various angles, as high as 90 deg, to the direction of heat flow. The colony spacing (C) is proportional to the square root of inverse freezihg rate. The relationship is C = 15.5(dfs/d?)-1/2 where C is in µ and 8 is in sec. The ratio of colony spacing to lamellar spacing is greater than 20.0 and increases with a decrease in the freezing rate. A duplex dendritic structure is produced at intermediate freezing rates. A fine lamellar eutectic is arranged within the dendrites (exhibiting side branches at an angle close to 60 deg from the main stem) and a coarse irregular eutectic appears in the interdendritic regions. The duplex eutectic structure is also produced at slow freezing rates. However, at slow freezing rates there is a Platelat of CuAl2, along the center of the main stem of each dendrite and the other lamellae are arranged perpendicular to the central platelet. THE eutectic between CuA12 and a! aluminum has been reported to freeze in a lamellar form by several workers.&apos;-3 chadwick4 has measured the interlamel-lar spacing as a function of growth rate. Kraft and Albright2 have reported on irregularities in the lamellar structures, and have proposed growth models which account for the formation of faults during solidification. In certain instances the lamellar eutectic has been found to exist in colonies. The colony formation315 has been attributed to the breakdown of a planar liquid-solid interface due to rejection of impurities. The aim of the present work is to study the structures produced from the eutectic aluminum-copper alloy under relatively fast solidification rates, such as encountered in casting and welding operations. The solid-liquid interface presumably remains planar under conditions of slow unidirectional freezing which produce lamellae aligned parallel to the direction of heat flow. The local growth velocities are the same over the entire interface and are equal to the rate of growth of the all-solid region. The spacing between the eutectic lamellae is inversely proportional to the square root of the growth rate of the all-solid region. Under the freezing conditions used in the present study, the solid-liquid interface is cellular or dendritic and the local growth velocities are different in the different regions of the interface. The relationship between the growth rate of the all solid region and the local growth velocities varies with the location and the shape of the interface. The growth rate of the all-solid region is, therefore, an inadequate parameter to describe the eutectic micro-structures which depend upon the local growth velocities. For this reason the structures have been examined as a function of freezing rate, dfs/d?, where fs is the fraction solidified at time 0. The freezing rate was varied by a factor of 4000. The relationship between the freezing rate, dfs/d?, and the growth velocit of the all solid region depends upon the specimen geometry and the shape of the interface. EXPERIMENTAL PROCEDURES The A1-33 pct Cu alloy used throughout this study was made in an induction furnace, using electrolytic copper and aluminum of commercial purity (99.7 pct), the primary impurities being silicon (0.12 pct), iron (0.14 pct), and zinc (0.02 pct). Three ranges of freezing rates were investigated: 1) A spectrum of rapid freezing rates (ranging from 0.40 to 7.30 sec-1) was obtained in arc deposits made on 2-in. thick cast plates of the eutectic alloy. The arc was operated at constant power and was made to travel at constant velocity on the surface of the plate that was in contact with the chill surface during solidification. The pool of liquid metal formed under the moving tungsten arc solidified rapidly by heat extraction through the unmelted plate. Conditions of unidirectional heat flow were achieved near the fusion zone interface, especially in the center of the arc deposits. The great advantage of the arc technique is that rapid cooling and freezing rates can be varied in a qualitative way. The correlation between the arc parameters and the solidification rate is given by the following relationship:6-8

Jan 1, 1970

By P. E. Doherty, B. Chalmers

Subboundaries may be revealed in aluminum by the formation of pits on the surface during cooling from elevated temperatures. The pits do not form in the vicinity of high- or low-angle boundaries. They are attributed to the condensation of vacancies from a super saturation produced during coolirzg. Using the vacancy pit and Schulz X-ray techniques for observing low-angle boundaries, a study was made of the transition from the nearly perfect seed to the striated structuke characterist-ic of aluminum crystals grown from the melt. It was found that the individual striation boundaries develop by the coalescence of very small-angle boundaries, as well as by the addition of individual dislocations. Several mechanisms for the formation of striations are discussed. Evidence was found suggesting that a super-saturation of vacancies exists near a growing interface, and it is proposed that the resulting climb of existing dislocalions produces "half&apos;-loops" at the interface, which combine to form the low-angle striation boundaries. LINEAGE, or "striation" boundaries, have been studied in detail by Teghtsoonian and Chalmers 1,2 in crystals of tin grown from the melt, and by Atwater and Chalmers3 in lead. They found that single crystals grown from the melt consist of regions which are separated by subboundaries that lie roughly parallel to the growth direction. A difference in orientation of 0.5 to 3 deg exists between the striated regions; the misorientation is such that the lattice of one region could be brought into coincidence with the lattice of its neighbor by a rotation about an axis approximately parallel to the direction of growth of the crystal. They observed an incubation distance for the formation of striations which increased with decreasing growth rate. They also found that in any crystal, the sum of all rotations of the lattice in one sense, in going from one striation to the next, is very nearly equal to the sum of all the rotations in the opposite sense. A striation boundary, which is a low-angle grain boundary, can be described as an array of dislocations. If it is assumed that suitable dislocations are introduced into the crystal during solidification, the formation of striation boundaries can be explained as a result of the migration of the disloca- tions into arrays. The formation of arrays is energetically favorable since the energy of an assembly of dislocations can be reduced by the interaction of the stress fields when a suitable array is formed. This investigation presents and interprets new information concerning the nature and origin of striation boundaries in aluminum. EXPERIMENTAL TECHNIQUE Single crystals of high-purity aluminum (Alcoa 99.992 pct) were prepared by horizontal growth from the melt.&apos;&apos; The specimens were subsequently electropolished in a solution of 5 parts methanol to 1 part perchloric acid kept between -10° and 0°C in a bath of dry ice and alcohol. The current density was approximately 6 amps per sq in. Doherty and Davis9 have shown that in aluminum sub-boundaries with misorientations of not less than several seconds of arc may be revealed by the vacancy pit technique. During cooling from elevated temperatures pits form on electropolished surfaces of aluminum crystals as a result of the condensation of vacancies.11 Pits do not form in the vicinity of small- or large-angle grain boundaries, presumably because such boundaries act as sinks for vacancies. Boundaries of misorientations down to 3 sec of arc are revealed as pit-free regions, see Fig. 1. The Schulz X-ray technique12 was used to determine the angular misorientations of subboundaries. In this method, white radiation from a micro-focus X-ray tube is used to produce an image of a fairly large area of a single crystal surface. Subboundaries cause splitting in the diffracted image, see Fig. 2. Misorientations down to about 15 sec of arc may be observed with this technique. OBSERVATIONS AND DISCUSSION Figure 1 shows a striated aluminum crystal grown at 10 cm per hr etched by the vacancy pit technique. An incubation distance of about 1 cm is observed before the initiation of striation boundaries. Fig. 2 is a Schulz X-ray photograph of a striated crystal similar to that shown in Fig. 1. A large area of the crystal was studied by means of a series of photographs. Fig. 2, which is a reflection from the (100) plane, included about the first 4 cm of crystal to freeze. There is an incubation distance of about 1 cm, and a distance of about 2 cm over which the angle of misorientation builds up to its final value of approximately one degree. Some twist component can be seen in Fig. 2 at the right side of the photograph. From Fig. 2 it can be seen that the sum of all rotations of the lattice in one

Jan 1, 1962

By J. B. Cohen, S. L. Sass

The nucleation process for y&apos;, Ni3Ti, is shou&apos;n to change from heterogeneous to uniform as the undercooling within the phase boundary increases. As unifornz nucleation beconres copious, (100) alignment is observed in the early stages of aging; howecer. whether or not the alignment is due to spinodal decomposition could not be determined. The presence oi a large-scale reversible segregation of titanium and not the intermediate precipitate y&apos; may be the cause of the initial strengthening in this alloy system and the discrete regions detected magnetically by Ben-Israel and Fine. Evidence is presented for the role of dislocations in the formation at long aging tines oJ the equilibrium precipitate . ThE nickel-rich portion of the Ni-Ti phase diagram consists of a terminal fcc solid solution of titanium in nickel called y.&apos; For temperatures below 1290°C if the solubility limit is exceeded and the temperature is high enough, the excess titanium will combine with nickel to form q, Ni3Ti, a four-layer hcp ordered structure with stacking sequence ABAC."~ Since a change in structure is involved, formation of q is slow and a metastable precipitate referred to as y&apos; forms. This metastable precipitate is an ordered cubic phase in the unconstrained state having a structure similar to Cu3Au (LIZ) with the composition Ni,Ti."&apos;5 From X-ray studies,8, 7 the following sequence is indicated for the precipitation reaction at low temperatures: 1) The appearance of satellites around low-index diffraction spots and splitting of high-index spots, indicating the formation of a slightly tetragonal phase in an imperfectly periodic arrangement. The tetragonal phase is y&apos; or the depleted matrix, depending on the volume fractions of matrix and precipitate (and hence, the alloy composition); tetragonality apparently arises from coherency strains. 2) The appearance of diffraction spots from the equilibrium second phase, q. Ben-Israel and Fine&apos;" used magnetic methods to measure the changes in matrix composition during aging, and to estimate the precipitate composition in a Ni-10.1 at. pct Ti alloy. They observed that the alloy, solution-treated at 1270°C and quenched, contained heterogeneities with discrete compositions which gradually vanished upon aging. It was suggested that y&apos; may initially form as a defect structure with several possible compositions deficient in titanium as compared to Ni3Ti; the composition Ni3Ti develops with prolonged aging. The initial composition was Ni6Ti which would require a very large deviation in stoichiometry for y&apos;. It was noted that at 700°C the precipitation reaction was 80 pct complete after 2 hr and that the volume fraction of second phase was 0.2 after 1 hr. Mihalisin and Decker&apos; suggested that the equilibrium precipitate, besides being nucleated at grain boundaries, could also form at stacking faults. ~errick" examined the formation of in Ni-Cr-Ti, and suggested that a collapsed vacancy cluster within y&apos; acted as a nucleus for intragranular q at high temperatures, 820" and 900°C. After long aging times at lower temperatures, 650" and 750°C, diffraction patterns taken from regions containing ribbons of stacking faults showed spots associated with q. Aging greatly enhances the mechanical properties of these alloys, even at high temperatures, and the metastable precipitate y&apos; is thought to be responsible.&apos; At 600°C the hardness of a Ni-10.1 at. pct Ti alloy increases rapidly during the first 20 hr of aging and continues to increase slowly thereafter.&apos;" At 700°C the hardness reaches a maximum at 1 to 2 hr and then decreases.l1 In this paper the kinetics of formation of y&apos; and the mechanism of transformation to in a Ni-Ti alloy are described. Evidence is presented that suggests that large-scale regions of titanium segregation are present initially and these may have an important influence on the initial strengthening of the alloy: these regions may be the heterogeneities detected by Ben-Israel and Fine. I) EXPERIMENTAL TECHNIQUES The Ni-10.3 at. pct Ti-0.6 at. pct A1 alloy was prepared by vacuum melting. Its chemical analysis is given in Table I. A 0.007-in.-thick strip was sandwiched between foils of the alloy and placed in an open Vycor tube with titanium chips (to get oxygen) and solution-treated for ; hr at 1150°C in a Globar Furnace under a flowing argon atmosphere. Following a quench into an ice-brine mixture at -5" to -2"C, a sample to be aged at temperatures up to and including 700°C was sealed in an evacuated Vycor capsule before being put into a fused salt bath. Aging above 700°C was carried out on bare samples in a salt bath. When removed from this bath and quenched into an ice-water mixture, the sample was normally coated with a thin layer of oxide easily removed with emery

Jan 1, 1970

By J. J. Gilman, T. A. Read

FOLLOWING the deformation 01 zinc monocrys-tals, sharply bent basal planes are observed near several types of inhomogeneities. Three of these in-homogeneities have characteristics which are quite regular so that they can be studied and analyzed. These are compressive kink bands, "deformation bands," and the inhomogeneities near end restraints. The present paper describes experiments in which "deformation bands" were artificially produced, and bend plane phenomena are discussed in terms of dislocation theory. Also, two new bend plane phenomena are described. The importance of bend plane phenomena in the deformation of crystals is not widely recognized. Many phenomena may be explained in a manner similar to the discussion in this paper. Jillson1 has pointed out that the "punching effect" in zinc is a bend plane phenomenon and is not caused by prismatic slip.&apos; Bowles3 has suggested that they may be involved in diffusionless phase changes. Cahn4 has discussed the role of bend plane formation in the polygonization of zinc. Experimental Work Tensile Kink Bands: Because of the geometrical similarity between "deformation bands" and "kink bands" (compare Fig. 1 of this paper with Fig. 1 of the paper by Hess and Barrett"), the band shown in Fig. 1 of this paper will be called a "tensile kink band," and that shown by Hess and Barrett will be called a "compressive kink band." It is felt that the term "deformation band" should be reserved for banded structures in polycrystalline materials such as iron." Tensile kink bands seem to form spontaneously in aluminum crystals deformed by tensile loading.7-10 In zinc and cadmium crystals they do not form in good, carefully loaded specimens.&apos;." However, tensile kink bands can be produced artificially in zinc crystals. The present authors did this by scratching one of the flat surfaces of triangular crystals transversely with a sharp needle. Natural tensile kink bands caused by inhomogeneities sometimes appeared in deformed crystals which were identical in appearance with the artificially produced ones. Zinc monocrystals were grown by the Bridgman method in graphite molds. Chemically pure zinc (99.999+ pct Zn) was used and the molds were sealed inside evacuated pyrex tubes during growth. The crystal cross sections were equilateral triangles with a typical base of 0.210 in. The artificial kink band shown in Fig. 1 is typical of tensile kink bands in zinc. The band lies between two bend planes which run from upper right to lower left and is inclined oppositely to the slip bands which are sharply bent at the two bend planes. The general form of the artificial tensile kink bands was independent of the scratch depth (1 to 5 mils deep) and also independent of which side was scratched. These variables did cause variations, however. Deep scratches produced more localized kink bands than light scratches. Also, if the angle between the slip plane traces and a transverse scratch varied appreciably among the three sides, then localization of the resulting kink bands also varied. Furthermore, if the slip direction lay nearly parallel to the scratched side, the band was more developed near the scratched side than at the opposite edge. Scratches produced tensile kink bands for crystal orientations from xo = 15" to x, = 75". Fig. 2 shows a scratched crystal after deformation. One triangular side lies in the plane of the photograph. The right hand tensile kink band was produced by a transverse scratch on the upper right side. The next two kink bands were the result of scratches on the front surface. The kink band at the left was caused by a scratch on the lower back side. All four bands have the same general form. A longitudinal scratch was also made on the crystal shown in Fig. 2 to determine the effect of a scratch on the critical shear stress. The critical shear stress of the scratched region was 33.9 g per sq mm compared to 24.4 g per sq mm for the un-scratched region above it. Fig. 3 shows Laue patterns of the crystal shown in Fig. 2. Fig. 3a shows the pattern of the undeformed crystal. The orientation was x, = 21°, A, = 31". After deformation, Fig. 3b was made of the homogeneously deformed portion of the crystal. The spots are compact but split into two halves. This region was elongated 45 pct and its orientation was x = 14", X = 20"; the sine law predicts x = 14", A = 20.5". Fig. 3c was taken near the center of the middle tensile kink band of Fig. 2. The pattern shows a range of orientations and polygonization in this region. The spread in orientation was due to the fact that the basal planes were curved (see Fig. 1) rather than flat as in the ideal case. Some may also have been the result of elastic distortions and "local curvatures." The orientation range was x = 23" to 32", A = 30" to 42". It is apparent from Fig. 3 that the material inside and outside the kink band rotated in opposite directions with respect to the tension axis during deformation. The orientation calculated from the ideal configuration of Fig. 9,

Jan 1, 1954

By W. J. Helfrich, R. A. Dodd

Binary aluminum solid-solution alloys containing various amounts of silver, magnesium, and zinc were prepared by careful directional solidification, and the hydrostatic and X-ray densities were compared. With the exception of the Al(Ag) alloys, the X-ray densities were consistently greater than the hydrostatic measurements, in agreement with earlier observations by Ellwood. In contrast to Ell-wood&apos;s interpretation in terms of vacant lattice sites associated with Brillouin zone effects, a tentative explanation based on the existence of solidification microshrinkage was favored. This hypothesis was confirmed by an examination of Al(Zn) alloys prepared by vapor diffusion of zinc into aluminum. The hydrostatic and X-ray densities were now in very close agreement, and it was concluded that the filling of Brillouin zones in aluminum solid-solution alloys does not necessarily result in the formation of defect structures containing an excess of vacant lattice sites. ThE existence of defect structures of the vacancy type in alloys in which the excess vacancies have an electronic rather than a thermal or mechanical, and so forth, origin is well recognized. Examples of incomplete lattices of this type are to be found in the Ni-Al,1-3 Fe-Ni-A1,4 c~-Ni-Al,5 Fe-Cu-Al,= and Co-A17 systems. These defect structures are of a special kind in that the intermediate phases possess an ordered atomic arrangement or superlattice, and in some instances the vacancy concentration may be unusually large, e.g., at 45.25 at. pet Ni in NiA1, approximately 8.8 pet of the lattice sites are unoccupied. Ellwood8-10 has reported similar defect structures in the aluminum solid solution alloys of the Al-Zn and A1-Mg systems and in alloys of the Au-Ni system." In Al(Zn) the (apparent) vacancy concentration rose, somewhat irregularly, to a maximum of about 2 pet vacant sites at 25 at. pet Zn, while in Al(Mg) the (apparent) vacancy concentration increased continuously to 1.7 pet at 15 at. pet Mg. An explanation in terms of Brillouin zone overlap was attempted, although Pearson12 has pointed out the difficulty of reconciling the observations with zone theory. However, the possibility of the effect being caused by the Fermi surface just touching a plane of energy discontinuity inside a prominent Brillouin zone has, in general, been accepted. In fact, Massal-ski13 has interpreted Ellwood&apos;s8 observations as confirmation of Leigh&apos;s14 theoretically predicted zone overlap occurring at approximately 2.67 electrons per atom. Unfortunately, Massalski was apparently unaware that Ellwood9 had revised his earlier results considerably, and the revised data did not confirm Leigh&apos;s analysis. Ellwood&apos;s clata were reexamined by the present authors who noted a possible correlation between the percentage defects as a function of alloy composition and the temperature interval of solidification measured from the respective equilibrium diagrams. This suggested an explanation in terms of shrinkage porosity rather than vacant lattice sites, and pointed to the desirability of reexamining appropriate alloy systems using: both Ellwood&apos;s method of specimen preparation (casting followed by wrought fabrication) and alternativ&apos;e methods, i.e., diffusion, which might be expected to minimize, or even completely obviate, microporosity. ALLOY PREPARATION 1) Cast Allolys and Aluminum Single Crystals. Al(Ag), Al(Mg;l, and Al(Zn) alloys of various compositions up to 20 at. pet silver, 13.5 at. pet mg, and 30 at. pet Zn were prepared by melting under helium and casting into graphite molds. In the first two systems, the maximum alloying addition was quite close to the limit of solid solubility, but the possibility of transformation to a&apos; during quenching somewhat restricted the suitable Al(Zn) composition range. The alloys were prepared from high-purity aluminum, a lot analysis showing 0.002 wt pet Cu, 0.002 wt pet Fe, and 99.996 wt pet A1 by difference. The silver, magnesium, and zinc were of 99.99+, 99.98+, and 99.998 wt pet respectively. Each composition was analyzed chemically. The as-cast ingots measured 7/16 in. diam and 5 in. length. One in. was removed from the top of the ingot, and the bottom 3 in. was machined to 0.275 in. diam; a point was also machined on the smaller diameter end. The remainder of the original ingot served as a top riser during subsequent remelting and controlled solidification. The machined ingots were now remelted using a Bridgman soft-mold technique to ensure directional solidification and, therefore, a minimum of micro-shrinkage. Alumina powder was used as mold material contained in an alundum thimble, and this crucible was placed in a helium-filled Vycor tube. The assembly was lowered through a suitable temperature gradient at approximately 0.5 in. min-l, and the risered portion of the casting was subsequently removed by sawing.

Jan 1, 1962

By T. Cantwell, N. C. Rasmussen, H. E. Hawkes

Beryl is a mineral that may be difficult to distinguish from quartz by casual field inspection. The easily recognized green color and hexagonal crystal form of coarse-grained beryl are by no means universal, even in beryl from pegmatitic deposits. If it occurred as a fine-grained accessory mineral in an igneous rock, it would almost certainly escape detection unless samples were submitted for petrographic or chemical analysis. There may be substantial deposits of some beryllium mineral, other than beryl, that has been overlooked because that mineral also closely resembles the common rock-forming minerals. A reliable and simple method of identifying beryllium minerals and determining the beryllium content of a rock would be helpful in exploration. This article describes preliminary experiments in applying nuclear reaction to the qualitative identification of beryl and to the semiquantitative determination of the beryllium content of rock samples. Gaudin,1,2 the first to apply a nuclear reaction in detecting beryllium minerals, developed a method that irradiates the sample with gamma rays, which react with beryllium nuclei to produce neutrons. The neutrons are then measured with standard equipment. The cross section for this reaction is about 1 millibarn. The cross section is a measure of the probability that a reaction will take place, for example, between a beryllium nucleus and an incident gamma ray or alpha particles.3-5 At 1-millibarn cross section for the reaction, satisfactory performance required a source strength of the order of 1 curie (3.7 x 10"&apos; disintegrations per sec, where each disintegration releases one or more gamma rays). The reactions will not take place if the gamma radiation is below a minimum energy, in this case 1.63 mev. The size of the source and the energy of the radiation made heavy shielding necessary for these experiments, both to reduce the background count of the neutron counter and to safeguard personnel. The original discovery of the neutron by Chad-wick in 1932 resulted from experiments with another nuclear reaction, induced by bombarding beryllium with alpha particles in which the products are carbon-12 and neutrons. The equation for this reaction is as follows:&apos; " ,Be" + ,He&apos;? 6C12 + 8,n&apos; [1] re-particle neutron In the above nuclear equation (Eq. 1), the sub- script number indicates the number of protons in the nucleus (the atomic number) and the superscript the total number of neutrons and protons (approximately the atomic mass). For the alpha-neutron reaction the cross section is about 250 milli-barns, or 250 times that of the gamma-neutron reaction used by Gaudin. The positively charged alpha particle is repelled by the positive charge of the beryllium nucleus; it must, therefore, have a certain minimum energy in order to approach close enough to the beryllium nucleus to react. For reaction with the beryllium nucleus, the lower limit of the alpha-particle energy is 3.7 mev. The alpha-neutron reaction, with polonium-210 as an alpha source, was selected for the present experiments. Alpha particles are emitted by polonium-210 at 5.30 mev, which is adequate for the reaction with beryllium. Furthermore, this isotope of polonium emits alpha particles with negligible associated gamma radiation, thus eliminating the necessity of shielding. The half-life of polonium-210 is 138 days. Inasmuch as alpha particles carry a possible charge and are large compared with most nuclear particles, their energy is rapidly dissipated in passing through matter. Their range in standard air is 3.66 cm,3 and they penetrate only a few tens of microns into a mineral sample. The short range in air can be minimized by preparation of a flat sample surface that can be brought very close to the alpha source during analysis. On the other hand, short range of alpha particles in air lessens the radiological health hazard and makes it possible to use this method without shielding. It must be emphasized, however, that the alpha emitters are potentially very dangerous if they enter the human body. Polonium must be handled with extreme caution. The literature has reported experiments on the yield of neutrons from reaction of alpha particles with beryllium nuclei. Feld" reports that in intimate mixtures of polonium and beryllium, 3 x 106 eutrons per sec are produced per curie of polonium. Elsewhere in the same reference it is stated that a sandwich-type source yields about one third as many neutrons as an intimate mixture. A table of neutron yields for full energy polonium alpha-particles on thick targets as reported by Anderson7 is the basis of Table I. From Table I it can be deduced that the elements most likely to interfere, i.e., those that also produce neutrons when bombarded by alpha particles, are boron and fluorine. These data also show that it will probably not be possible to determine very small quantities of beryllium in rocks because of the masking effects of the major elements, sodium, magnesium, and aluminum. The neutrons emitted in the alpha reaction are detected by another nuclear reaction. Either of the

Jan 1, 1960