Search Documents

By George Hallerman, R. J. Gray

In the customary method of studying age hardening, the process of aging is interrupted by cooling the specimen and measuring its room-temperature hardness. However, the aging process may be conveniently measured while it is occurring at an elevated temperature by a series of hardness measurements at suitable time intervals. The latter method gives data void of complications that might result from processes occurring during the quench, and the data may be more indicative of the alloy's properties at elevated temperatures. The elevated-temperature hardness tester1 used in the present study employed a dead-weight loading system with a synthetic sapphire Vickers-type in-denter mounted in a molybdenum holder. The tests were performed in a vacuum of l0 to 10 torr. Hardness impressions were made by applying a 2-kg load for 15 sec. Although the hardness impressions made in the tester at elevated temperatures were measured at room temperature, the correction for thermal contraction during cooling was insignificant and therefore neglected. The two methods of obtaining age-hardening data were compared using eighteen samples of Haynes alloy No. 25* which had been annealed for 2 hr at 2250°F (1232°C). The test surfaces of the 1 by 1 by 1/4-in. samples were ground and then polished on vibratory polishers2 prior to testing. Three groups of five samples were aged at 1400°, 1650°, and 1800°F (760°, 899", and 982°C) for 4, 16, 32, 64, and 100 hr and then water-quenched. Room-temperature hardness as a function of aging time for these samples is shown in the broken lines of Fig. 1. Only one sample for each of the three temperatures was necessary to determine the hardness of the alloy at the aging temperatures with the elevated hardness tester. The specimens were heated isothermally and hardness impressions were made after various time intervals up to 100 hr. The results of these determinations were represented by solid lines in Fig. 1. Comparison of the two sets of curves shows, as expected, the hardness values determined at the aging temperatures are lower than the room-temperature hardness of the quenched specimens and the hot-hardness values reflect a strong tempera-

Jan 1, 1964

By J. Murray Riddell

The method of treating hazards wherein value is decreased, is cited by R. D. Parks. Quite properly, the theory of probabilities is made use of when multiple hazards are under consideration. E. F. Fitzhugh, Jr. $ deals with the enhancement phase graphically and uses percentage of expectancy as derived by a study of conditions and the selection of a chance factor. It is believed that both of the authors referred to fully recognize the part which human judgment plays, and as coupled with experience, training, and the application of the theory of probabilities; in that respect the composer of this paper is in accord. Within the past year, the task of arriving at a mathematical method for handling the enhancement phase, and, additionally, coupling the negative and positive phases, was explored. In citing the hazard, or negative phase, reference is made to the text by Parks. Types of Items far Consideration On the hazard, or negative, side some of the common items to be considered are: 1. Subsidence and resulting damage. 2. Flo wage of water above that anticipated. 3. Geological. (a) Decline in quality of probable ore, and/or (b) Decline in quantity of probable ore. 4. Depressed economic conditions. 5. Litigation. Similarly, on the enhancement, or positive, side common items to be considered are: 1. Flo wage of water below that estimated. 2. Geological. (a) Increase in quality of probable ore, and/or (b) Increase in quantity of probable ore. 3. Elevated economic conditions. 4. Litigation. 5. Technological advance in the art. Tn all of the above items no consideration is given to placing a value on, or giving consideration to, possibilities. For example, referring to Table 1, the estimated normal present value (Vp) of a property, exclusive of risk consideration, is taken as $100,000. If subsidence affects the usefulness of the shaft, the present value of the property will decrease to $60,000, and it is judged that there is a 1 in 4 chance that the shaft will be impaired;, consequent.ly, if subsidence took place, the present value would be reduced ($100,000-$60,000) $40,000, but since the chance is 1 in 4, the allowable reduction would be only one-fourth of $40,000, or $10,000, or a discount of 10 pct (hazard factor) of the original value; the related safety factor is taken as unity minus 0.10, or 0.90. Thus, the original Vp, $100,000, multiplied by the safety factor ($100,000 X 0.90) gives $90,000, which is the adjusted present value of the property in this single instance. Additionally, there is a water hazard, which, if encountered, will reduce the present value to $85,000. The chance is 1 in 3. Following through with the same reasoning as shown above, the derived safety factor will be 0.95. To arrive at a single multiplying factor which will express the combined risks, the two safety factors are multiplied together, 0.90 X 0.95 = 0.855. That factor when multiplied by the origirial Vp will give the adjusted, and

Jan 1, 1950

By George Ashley

IT has been said that the Pittsburgh bed is the most valuable single mineral deposit yet known to man. The figures in Table 1 are presented in substantiation of that claim. TABLE 1.-Pittsburgh Coal Bed to Dec. 31, 1934 Estimated Recov- State Thousands of ( Value erable Reserves Tons Mined (1935), Thousands of Tons Pennsylvania 3,144,000 $4,922,000 7,500,000 Ohio 493,391 843,699 3,000,000 West Virginia 650,000 1,027,000 5,500,000 Maryland 183,224 254,681 20,000 Totals 4,470,615 $7,047,380 16,020,000 PRODUCTION AND VALUE In order that the figures of Table 1 may have any weight, a statement must be made to show how they were arrived at. Because, to the writer's knowledge, no accurate separate figures ever have been kept of the production and value of the Pittsburgh coal bed, the states will be considered separately.

Jan 1, 1936

By Victor V. Damiano

THOMPSON1 proposed a very convenient nomenclature in which an imaginary regular tetrahedron was used to facilitate the analysis of dislocation interactions in fcc materials. The faces of the tetrahedron were constructed parallel to four different sets of (111) planes, and the six edges of the tetrahedron then represented the six 1/2 [110] Burgers vectors. The lines which joined the midpoint of a tetrahedron face with the ends of one of its edges represented the Burgers vectors of the partial dislocations into which the six 1/2 [I101 Burgers vectors dissociate. Using Frank&apos;s2 notation, all of the dislocations and their partial dislocations were identified and numerous possible interactions were postulated. The case of hexagonal crystals is unfortunately more complex, but it appears convenient to introduce a system similar to that of Thompson1 to describe the various slip planes and Burgers vectors encountered in cph structures. The system proposed has limited application but serves to identify the following slip systems {0001} <l20>, {1010} <1120>,

Jan 1, 1963

By H. Doi, Y. Fujiwara, K. Miyake

In order to throw light on the mechanism of plastic deformation of WC-Co alloys, compressive tests of WC-(7 to 43) vol pct Co alloys have been carried out at room temperature, and stress-micro strain relation has been investigated in detail. The analysis of the factors affecting the yield stresses reveals that the yield stresses can be predicted by modified Oro-wan&apos;s theory if one properly estimates the planar in-terfiarticle spacings. Conzpressive straining of some of the alloys by 0.066 to 0.17pct increases the decrements by a factor of as much as 3.4 to 14, whereas the corresponding increase in the electrical resistivities is less than 10 pct. The analysis of the decrement data in terms of -Gramto and Lücke theory shows that the marked increase is attributed to increased dislocation darnping itt the binder (cobalt) phase. By cornbilling the decrement data and the conzjwession duta, one obtains the relation between flow stress in shear (?t) and increase in dislocation density (p): At = const . v6 . This is interHeted to mean that the mechanism of strain hardening of CirC-Co alloys is essentially sarne as the one for dispersion strengthened alloys. The possible effect of bridge formations between the carbide particles has also been examined. OWING to the combination of hardness, strength, and other physical and chemical properties, WC-Co alloys have opened the way for unique fields of applications, the recent ones being, for instance, anvils for super-high-pressure generation apparatuses. In such applications, the alloys are frequently subjected to very high compressive stresses: these stresses may cause the alloys to deform plastically and eventually to fail. However, much remains obscure regarding the nature of the plasticity of the alloys. Evidently, the alloys owe their high strength to the hard carbide particles which frequently occupy as much as 80 to 90 pct in volume fraction, whereas the ductility required for practical applications is provided by the small amount of the binder phase between the carbide particles. When the volume fraction of the carbide phase is not very large, deformation behavior of the alloys may be described by some of the current dispersion strengthening theories. However, greatly increasing the carbide phase is thought to lead to some carbide skeleton structure or bridge formations owing to the increased chances for direct contacts between the carbide particles;1,2 this may appreciably affect the plasticity of the alloys. Regarding the effect of formation of the carbide skeleton structure, it is interesting to note the work by Ivensen et al.3 on compression tests of the alloys containing somewhat large carbide particles; they observe extensive generation of slip bands in the carbide particles after application of some preliminary compressive stresses. They interpret the results in terms of plastic deformatiot: of the carbide particles which are supposed to have formed a skeleton structure; the binder phase plays only a passive role, at least in the early stages of the deformation. That carbide crystals exhibit microplasticity at room temperature is apparent from the work of Takahashi and Freise4 and French and Thomas5 on indentation of WC single crystals. On the other hand, Dawihl and coworkers6-10 maintain that even when volume fraction of the carbide phase is very large (for instance, more than 90 pet), a very thin binder layer generally exists between the carbide particles. They interpret the results of the extensive mechanical tests in terms of the plasticity of such a layer. Gurland and Bardzil11 point out that decrease in ductility of the alloys with increase in the carbide phase is caused by the effect of plastic constraint exerted by the dispersed carbide particles. Drucker12 further develops this concept from a continuum-mechanics approach on an assumption that a continuous thin binder layer separates the carbide particles. A common feature of the studies reported so far on the plasticity of the alloys is that the information deduced is invariably qualitative in nature. Thus, very few systematic experiments for obtaining reliable and sufficiently detailed stress-strain curves of the alloys varying widely in the microstructural features have been carried out. In particular, it may be of special interest to investigate in detail the early stages of the plastic deformation of the alloys in order to shed light on the strengthening mechanism. However, such work appears to be extremely rare. Doi et al.13 recently reported a first brief account of the results of some quantitative analysis of the plasticity of the alloys in terms of dislocation theory. Their experiment was rather limited in the composition range covered (volume fraction occupied by the carbide phase: 79 to 83 pct), and thus they could not necessarily elucidate the controlling mechanism of plastic deformation of the alloys of a more general composition range. Consequently, in the present investigation, deformation behavior and some other physical properties of the alloys were investigated and discussed in more detail over a much wider composition range. SPECIMEN PREPARATION WC-Co alloys used in this experiment were prepared in cylindrical or rectangular form by sintering in vacuo compressed mixtures of tungsten carbide and cobalt

Jan 1, 1970

By D. F. Haley

THE molybdenite deposits at Climax, Cool., have recently attracted considerable notice, because of their great size, as compared with other known deposits of the same mineral. Climax station, on the. Colorado Southern narrow-gage railway, was at one time the Fremont Pass station on the Denver & Rio Grande. It is 1412 miles northeast of Leadville, and exactly on the divide between the headwaters of the Arkansas and Blue Rivers. Climax station is 11,300 ft. (3444 m.) above sea level. The working tunnel of the molybdenum mine is about 1 mile from the railroad, and 900 ft. above it, at an elevation of 12,200 feet. The climatic conditions are severe, and snow can be expected during any month of the year. Winter sets in not later than Nov. 1, the snowfall is very heavy, and the ground remains covered with snow until June 1, and sometimes even later. From June 1 to Oct. 1, there is much rain and a little snow falls in the late afternoon nearly every day. In October, 1917, there was one of the biggest snowstorms and blizzards of the year. Being just on the divide, constant strong winds blow the fine, dry snow with such force that, during most of the winter, outside work of any kind -is accomplished under most trying conditions. Within a few miles of Climax, some of the richest gold placers of Colorado occur. One of them, McNulty Gulch, was just across on the northeastern flank of Bartlett mountain. This mountain, therefore; was well prospected, and the presence of a bluish mineral and a yellowish one resembling sulphur was known to the old prospectors. They thought for many years that the blue mineral was graphite, and it is difficult to say when it first became definitely known that it was molybdenite. Claims had been staked as .early as 1896, but it is evident that the prospectors at that time were looking for gold. In 1902; several claims were acquired by H. Leal, who drove a tunnel about 900 ft. (274 in.) long, probably to intersect a fault, which shows on the surface, and was thought to be a possible gold carrier. Gold was not found in paying quantities, but throughout its length the tunnel passed through a fractured silicified granite, showing a fairly uniform distribution of a fine-grained bluish mineral, which was finally determined as molybdenite. For several

Jan 8, 1918

By B. Ramaswami, Y. Nakada, R. Schaefer

EARLY investigations1 of the tensile properties of fcc single crystals did not reveal the easy glide region of crystals with axial orientations inside the standard stereographic triangle. With the production of high-purity metal, a region of easy glide became known1 Consequently, modern work of fcc single crystals ordinarily uses high-purity metal. In order to determine whether further purification of aluminum beyond 99.992 pct purity would result in significant alteration of the tensile properties of single-crystal specimens, metal was zone-refined and crystals grown from the purified material were compared to crystals grown from the original material. Some similar studies were made of gold and silver crystals. Aluminum. The material to be zone-refined was cast into bars 1/4 by 1/2 by 8 in. A zone, approximately 1-1/2 in. in length, was passed over each bar eighteen times at a rate of 6 in. per hr. Melting and stirring were carried out by an induction heater under a vacuum of approximately 10"5 mm Hg, the metal being held in a graphite container. he last 3 in. of each processed bar were then cut off and discarded. No quantitative determination of purity was made, but the effectiveness of the zone refining was tested by a method first suggested by Walton, Tiller, Rutter, and winegard3 depending on the relation between cell structure and impurity content. Rutter and Chalmers4 have shown that a cellular substructure is associated with the presence of impurities during solidificatibn, and have shown how solute rejection at the freezing interface can lead to such an association. Thus the observation of cells gives evidence of impurities being present. To look for cells, the last quarter of a zone-refined bar was cut off, remelted, and grown rapidly into a single crystal. It was then heavily electropolished and thermally etched.~ Cells were revealed by the presence of many parallel rows of etch pits. The same process was carried out with material from the first half of the bar, but no cells were detected. These observations give evidence that a significant amount of impurity material was moved to the "dirty" end of the zone-refined bars. Further indication of purification was obtained from recrystallization experiments. It was found that, after cold rolling to a 97 pct reduction in thickness, the material from the pure end of the bar re-crystallized at room temperature, but the material from the "dirty" end did not. Other investigators6 have previously reported that zone-refined aluminum recrystallized at room temperature. The purified material was cast into 1/4 by 1/4 in. bars, and seeded crystals were grown in a (100) orientation and in a single slip orientation where the Schmid factor is m =2. The crystals were then cut to the desired length, annealed for 24 hr at 610°C, furnace-cooled, and etched in Tucker&apos;s etch. The crystals were tested in a floor-model Instron machine, using a gage length of 4 in. and a strain rate of 1/2 pct per min. The resulting curves are shown in Fig. 1. It is seen that the differences between the curves for zone-refined and original-purity materials are small. These differences are within the scatter observed for crystals of the original purity. In particular, the zone refining does not significantly affect the onset of the different "stages" of the curves.

Jan 1, 1964

By W. J. Slatosky

W. 0. Philbrook (Cairiegie Institute of Technologyogv—Mr. Slatosky has presented an interesting and constructive paper that represents another step along the way of converting steelmaking from an art to a science. I am confident that his computer will be practical and successful and that with a very few months of experience it will provide a significantly better degree of control than his record of 65 pct of heats within range obtained with the slide-rule calculator . A paper such as this, with a lot of symbols and condensed mathematics, is difficult to comprehend quickly. Since I have had an opportunity to study it carefully, perhaps my evaluation of its validity and accomplishments will save time for others. Mr. Slatosky has correctly used standard principles of stoichiometry and heat balances, which are available to anybody, but he has also brought to them two original contributions: 1) He has developed from operating data some empirical relations for predicting the final FeO content of the slag (at 0.5 pct C end-point) as a function of slag basicity, lance height, and scrap, ore, and scale in the charge. This improves the accuracy of prediction of temperature or scrap requirement compared with assuming an arbitrary, constant FeO content at the end of each heat. There is no assurance yet that exactly the same relations will hold for other furnaces or practices, but similar correlations can be expected. 2) He has combined calculations that are ordinarily carried out laboriously as a number of individual steps into a single, simple linear equation that can readily be fed into a machine. This involved a tremendous amount of painstaking detail work as well as the imagination to see the possibility and work out the steps. While his particular Eqs. [3] and [6] are valid only for the furnace design, charge weight, and blowing time used at Aliquippa Works, only a few numerical values have to be changed to adapt it for other conditions. In order to arrive at a useable solution, Mr. Slatosky had to make some basic assumptions about the process that are similar to those used by others. He neglected variation in some process variables and assumed an arbitrary average value for waste gas analysis and temperature for want of more exact information at the present time. All of these judgments are clearly stated. In addition, some thermody-namic data presently available are not adequate for the job, notably in relation to heats of formation and sensible heat in slag, and some expedient has to be adopted to get around the difficulty. Other people might prefer slightly different judgments about these details and hence obtain somewhat different numerical solutions. This is not of serious importance, however, because the errors accumulate in the "heat loss" term and are largely self-compensating for a constant heat time. Although the extended Eq. l(a) in Appendix I was set up as a rate equation originally, for convenience in using an analogue computer as stated in the paper, the time dependence was removed by later mathematical manipulations and assumptions about the process. The final result is an integration of element and energy balances from initial to final states; this procedure is as legitimate here as in any other form of heat-balance calculation. The formal handling of stoichiometry and thermochemistry appears to be correct, and it is assumed that any arithmetical errors would have come to light in applying the calculations to furnace practice. Mr. Slatosky&apos;s approach is not necessarily unique, in that other people might start with apparently different equations or prefer another form of final equation for another type of computer. However, he has presented an accomplished result that appears to be a theoretically sound and practically useful way of applying scientific principles and rapid computation for better control of steelmaking. His success will undoubtedly encourage himself and others to improve on the mathematical model and its use as better informatioq becomes available. John F. Elliott (Massachusetts Institute of Teck-t2ology)-The last comment by Mr. Richards that a calculator is quite unnecessary for an L-D operation ?-equi??es a rebuttal. The L-D furnace is a very high capacity process which places a premium on close control. When one is making steel at rates between 100 and 200 tons per hr, one cannot afford the luxury of an extra 5 or 10 min at the end of a heat correcting for an error that should never have been made in the first place. Mr. Slatosky&apos;s paper is a very sound application of the simple principles of stoichiometry and the energy balance. It is a satisfactory and valuable start, but only the start of the development of methods of control for this process. An analysis of the process shows that it should be very suitable to control by a computer. This is especially the case when various grades of steel are to be made. In fact, it would seem that the organizations who are planning new and bigger installations of L-D vessels should consider carefully the advantages that would stem from computer control of a vessel with the operator present to do little more

Jan 1, 1962

By H. K. Martin

IN September 1954 the Lavender pit concentrator at Lowell, Ariz., began treating low grade porphyry copper ore from the nearby Lavender mine. Nominal capacity of the mill is 12,000 tpd, but production now averages 16,226 tpd. Plant design was somewhat affected by the re-use of old equipment and the substantial re-use of old buildings. A breaking-in period between June 24 and Sept. 1, 1954, was devoted to training operators and making minor changes in launders, chutes, and piping. From Sept. 1, 1954, to July 1, 1956, the mill treated 8,598,105 dry tons of ore, from which 556,541 tons of concentrate were produced, containing an indicated 144,055,185 lb of copper.

Jan 11, 1957

By P Shenon

THE Elk City district is in north-central Idaho about 60 miles east of Grangeville and near the headwaters of the South Fork of the Clearwater River (Fig. 1). At the height of its boom in the early sixties, Elk City had a population of more than 2000 people.1 Published estimates of the placer production from the district range from $10,000,000 to $18,500,000. Although the first quartz-vein location was made as early as 1870, very little gold was extracted from quartz ores until after the completion of the American Eagle mill in 1902. Between 1902 and 1932 a total of from $750,000 to $1,000,000 was taken from about six properties.

Jan 1, 1934

By R. W. Armstrong, P. C. Jindal

WE should like to point out that Hu and cline.1 in their study of the recrystallization of titanium, have obtained data which appear very important for accurately determining the dependence of hardness on grain size for this material. It has been proposed2,3 that the macrohardness. H. of a poly crystalline material should depend on its average grain diameter, l. in accordance with the relationship: H =H0 + kHl-1/2 [1] where Ho and kH are experimental constants. This relationship follows from the hardness dependence on yield stress or flow stress4 and. thence. the Hall-Petch stress-grain size relations.5&apos;6 Fig. 1 shows the Diamond Pyramid Hardness. Dph. values vs the reciprocal square root of the average diameter of the recrystallized grains, taken from the titanium data reported by Hu and Cline. The data in this figure give a linear relationship for the hardness over a wider range of grain diameters than anyone has previously measured. The data are not linear, for example. on an I-1 or l-1/3 scale. At the very large grain sizes the data are in reasonable agreement with measurements earlier reported by Holden, Ogden, and Jaffee.7 For the grain sizes smaller than 1 µ, which correspond in Fig. 1 to hardness values greater than 150, the material was reported as not being completely recrystallized and, therefore, these data appear less meaningful as far as the hardness dependence on grain size is concerned. However. it has been suggested by other investigators that deformed materials exhibit strengthening due to an effective grain size refinement even in the unrecrys-tallized condition.8-10 Also. Hu and Cline have shown that a change in texture occurs for titanium during the recrystallization process and. especially. following it. The influence of this texture change on the hardness-grain size relation would be expected to be less than that reflected in the unidirectional stress-strain behavior, even for an hcp material. because of the actual combined stress system which applies for the hardness test. For, at least, certain combined stresses, average values of the orientation factors should still occur in the stress-grain size analysis of a material exhibiting a preferred orientation.&apos; We previously suggested3 that the experimental constants. H0 and kH. from the Meyer hardness-grain size relation should be correlated with the Hall-Petch constants, uoy and ky.5,6 Our recent study" of the Hall-Petch relation and the hardness-grain size results which obtain for a number of polycrystalline materials shows that: The inequality [2] is explained by noting that the hardness should be correlated. perhaps, with the flow stress at some value of strain, as suggested by Dugdale,12 and. also, that the influence of strain on the Hall-Petch relation is to increase uo and not change k very much." Normally. Dph values and Brinell hardness values are nearly equivalent to each other and to the Meyer hardness number. For the case of titanium, then. the inequality [2] may be evaluated using the experimental values of ao = 8.0 kg per sq mm and ky = 1.3 kg per mm3/2 obtained from the experiment of Guard.6,13 Thus, (Ho/uo) = 12 and (kH/k3) = 1.3. This research has been supported at the University of Maryland through the Center of Materials Research by the Advanced Research Projects Agency of the U.S. Government.

Jan 1, 1969

By G. Ansell, J. Weertman

G. Ansell and J. Weertman (Rensselaer Polytechnic Institute and Northwestern University, respectively) —The great increase in creep strength that you found in your indium-glass composites is quite striking. You ruled out a decrease in the dislocation source density as an explanation of the reduced creep rates. Could you tell in greater detail why you feel this is an inadequate explanation of the phenomenon. It appears to us that it is a reasonable one. Consider the spacing between the glass spheres in your composites. It is of the order of 10-3 cm. The spacing between dislocations in a (single-phase) annealed metal with a dislocation density of the order of 106 to 107 cm per cm3 also is of the order of 10-3 cm. Therefore it is not reasonable to expect that the arrangement of the dislocation lines is the same in the composite material as it is in pure indium. If the dislocation arrangements are different the density of the dislocation sources likewise will be different. Another way of saying the same thing is the following. The presence of the glass spheres reduces one of the dimensions of the metal phase to 10-3 cm. Thus the metal phase, in effect, is equivalent to a metal whisker or a thin metal film. These are strong because they have so few dislocation sources within them. T. D. Gulden and J. C. Shyne (authors&apos; reply)—It is to be expected that the arrangement of dislocations in the indium-glass composites would be quite different than that observed in annealed pure metals. The dislocations probably exist primarily as segments connecting the dispersed particles. These segments and discontinuities at the very large area of particle-metal interface could act as dislocation sources. For this reason it was considered unlikely that the very high observed creep strengths were due to a low density of dislocation sources. A more likely explanation is that the distance a dislocation can glide is limited by the interparticle spacing. The particles averaged about twenty in diam and were less than 10 µ apart in the 40 pet composite. It seems reasonable to assume that a dislocation loop in a volume of metal surrounded by these relatively large particles would not bow out around them but would be restricted to the region between particles. A general expression for creep rate is given by the product of the dislocation source density, the average rate at which loops are generated from a source, and the average area swept out by a loop times the Burgers vector magnitude. Assuming that the area swept out is limited by the interparticle spacing and the rate of generation of loops is controlled by the rate of climb of loops over second-phase particles, as suggested by Ansel and Weert-men,9 a calculation of the dislocation source density may be made from the experimentally observed creep rates. The result comes out 10" to 10" per cm3 which seems reasonable in view of the assumptions involved. Thus it is concluded that the anomalously low creep rates that were observed were due not to a low dislocation source density but rather to a limitation on the area swept out by each dislocation loop.

Jan 1, 1964

By Milton H. Fies

THE Alabama Power Co. operates its Gorgas mine in Walker County, Ala., immediately adjacent to its steam plant. The coal mine property consists of some 19,000 acres of fee and mineral ownership lands, and all of the production from the mine-approximately 2500 tons per day-is utilized at the company&apos;s adjoining steam plant. The company is presently considering the purchase of additional acreage adjoining its present holdings which acquisition will almost double its reserves. The coal is transported from the mine over a short railroad spur in the company&apos;s own equipment, either directly to one of the large plants at Gorgas or to a storage pile. The longest haul is about 1 1/4 miles.

Jan 12, 1950

By W. D. Robertson, D. A. Koss, A. J. Goldman

THE significance of the hcp (E) structure, which appears when Fe-Cr-Ni alloys (stainless steels) are transformed martensitically, has been the subject of considerable study and speculation.&apos;-7 It has been proposed that the hcp structure is an intermediate phase in the transformation: fcc (austenite) — hcp (E) -bcc (0, martensite). However, recent studies by transmission electron microscopy6,7 have indicated that the hcp structure may be a consequence of the large shear strain (10 deg) associated with the fcc to bcc transformation. That is, the transformation to a bcc structure is accompanied by extensive dissociation of dislocations in the parent austenite, which has a low stacking-fault energy.&apos;-&apos;&apos; X-ray diffraction analysis of the transformed alloy reveals a pattern of lines corresponding to an hcp structure, together with the parent austenite and the bcc product. The fact that the hcp structure is observed only in association with the bcc product in these alloys1&apos;4&apos;6&apos;7 lends support to the conclusion that the hcp structure is a secondary product of the transformation. But, the evidence is circumstantial and more definite resolution of the two possibilities mains to be obtained. If the hcp structure is an intermediate product of the transformation, as in 1) above, then two distinct Ms temperatures should be observed, Ms and To evaluate this possibility, three different techniques that are particularly sensitive to the structural changes accompanying such transformations have been used in an attempt to detect two transformation temperatures in an Fe-15.1 Cr-11.7 Ni alloy: 1) high-sensitivity differential thermal analysis, 2) the temperature dependence of resonant frequency (elastic moduli), and 3) internal friction of cylindrical specimens vibrating in either longitudinal or torsional modes. To obtain the necessary sensitivity, a differential thermal analysis technique was developed for the purpose, Fig. 1. The differential temperature between a nickel standard and the specimen was amplified to produce a sensitivity of 0.02OC (0.7 v) in temperature difference. The specimen and the standard were surrounded by a copper block and the entire assembly was immersed in a solution of methyl and ethyl alcohol, freezing point, - 140OC, which provided the necessary thermal stability to reduce temperature fluctuations to negligible proportions. A cooling rate of about 1.5OC per min was obtained by bubbling liquid nitrogen through the liquid bath. In order to calibrate the experiment with respect to the evolution of heat in the transformation from fcc to hcp structures, differential cooling curves were obtained for an Fe-19.8 Mn alloy, which transforms martensitically to E.&apos;"-&apos; a Ms(e) temperature of 145°C was clearly defined by a temperature

Jan 1, 1964

By K. P. Gupta

In ternary systems containing silicon and transition elements, silicon was found to stabilize the a phase.1 At 1000°C a broad a-phase region was found in the Cr-Mn-Si system,1 which extends very close to the Mn-Si binary line (within about 3 at. pct Cr). Later investigation showed that, if chromium was replaced by small amounts (3 to 5 at. pct) of v,1 Ti, Nb, Ta, Mo, or w,1 the a phase or the R phase was again stabilized. Whether stabilization of complex crystal structures is possible by small additions of iron, cobalt, or nickel to Mn-Si is not known. Investigation of the Co-Mn-Si system4 shows the presence of several ternary phases, but at cobalt concentrations higher than 10 at. pct. In the present work the ternary systems Fe-Mn-Si and Co-Mn-Si were investigated at low iron and cobalt concentrations (< 7 at. pct Fe or Co) in search of phases of complex crystal structure. Experimental Procedure. The alloys were pre-pared by induction melting Co (99.8 pct), Si (99.9 pct), electrolytic Fe (99.88 pct), and Mn (99.9 pct) in recrystallized alumina crucibles under an inert atmosphere of argon gas. The alloy specimens were wrapped in molybdenum foil, sealed in argon-filled (at 1/2 atm) silica capsules and annealed for 72 hr at 1000" 1°C, and water-quenched. The use of argon-filled capsules was found to reduce the losses during annealing and the molybdenum foil prevented reaction of the manganese with silica. The total weight loss due to annealing and melting was found to be less than 1 pct for each alloy. The alloys were not analyzed chemically. A part of each annealed specimen was polished and examined metal-lographically using etching solutions, a) 0.1 to 0.5 pct oxalic acid solution in water, or b) 0.5 pct HNOs in methyl alcohol. Even though the etching solutions were quite dilute, the specimens were heavily etched and stained if the etching time exceeded more than 5 sec. Etching times found suitable were usually 1 to 3 sec. For X-ray diffraction work, a part of each annealed specimen was powdered by crushing after the exposed surface was ground off. X-ray diffraction patterns were taken with powder specimens using a 20-cm-diam asymmetrical focusing camera and Fe K radiation. These powder patterns were found quite satisfactory for the purpose of phase identification. However, X-ray diffraction patterns of single-phase alloy powder specimens were also taken using a Siemens diffractometer and filtered Cr-Ka radiation for greater resolution of the diffraction lines. Results and Discussion. In the Fe-Mn-Si system a new ternary phase was found at 1000°C between the p manganese and the MnsSi phase regions. The approximate phase relationships in this region of the 1000°C isothermal section are shown in Fig. 1. The new ternary phase, designated as the K phase, was found to occupy a narrow region about 1 at. pct in width and extending from Fee.oMn74.5Silg.5 to Fe.Mn.Silg.. At the high-iron content region, the K phase was found in equilibrium with the p manganese and the MnsSi phases but, at the low-iron content end of the K-phase region, the phases in equilibrium with the K phase were the p manganese phase, the MnsSi phase, and another new phase, designated in Fig. 1 as the v phase. The v

Jan 1, 1964

By D. H. Baldwin, R. E. Reed-Hill

WITH the polarized-light microscope, the basal-plane trace may be located on hcp crystal surfaces to an accuracy of approximately 1 deg. This determination can also be made on very small crystalline areas, approaching a micron in width. Measured basal-plane trace distributions on selected specimen surfaces can yield1,2 valuable texture information. However, the polarized-light microscope is probably best suited to orienting

Jan 1, 1965

By Paul A. Farrar, Louis P. Stone, Harold Margolin

The Ti-Mo-0 system was investigated in the region 0 to 45 wt pet Mo and 0 to 10 wt pet 0, from 600° to 1400°C. Solidus data are also presented. Isothermal sections at 700°, 900°, 1100°, and 1300°C, and vertical sections at 1 and 28 wt pet Mo, and 1 and 10 wt pet 0 are included. STUDY of the Ti-Mo-0 system has been carried out as part of a general study of titanium phase diagrams. The binary Ti-Mo system has been investigated in whole or in part by several investigators.1-8 It was found that molybdenum and ß-tita-nium form a complete series of solid solutions and that the ß/a + ß transus is depressed to below 600°C at approximately 30 pet Mo.* The Ti-O system has been investigated partially or fully in the range 0 to 30 pet 0 by several investigators." Extensive solubility of oxygen occurs in the field. While a is

Jan 1, 1957

By J. C. Coughlin, M. C. Flemings, T. Z. Kattamis

Experiments on Al-4.5 pct Cu alloy shou the dendrite arnz spacing increases with time during isothermal holding in the liquid-solid region. This coarsening of the dendrite structure is then shown to occur at roughly the same rate during solidification as during isothermal holding. Final dendrite arm spacing in both types of experiments depends primarily on coarsening kinetics. It is concluded that driving force for the coarsening is reduction of liquid-solid sur/ace area. Calculations based on two simple models show qualitative apreement with experinzent. 1 HERE has been considerable interest in recent years in determining effects of solidification variables on dendrite arm spacing in cast alloys. The interest has resulted primarily from the clear evidence that 1) the degree of homogeneity which can be obtained in a cast alloy with commercially practical homogeni- zation treatments depends on dendrite arm spacing, and 2) mechanical properties of the heat-treated casting depend on the homogeneity obtained.&apos;-&apos; In wrought alloys, microsegregation from the original ingot is carried to some extent into the final wrought product, and affects mechanical properties in similar fashion Experimental evidence available indicates that dendrite arm spacing in cast aluminum alloys is influenced only by cooling rate (or "local solidification timev*) and alloy content.Fig. 1, for example,

Jan 1, 1968

By J. P. Hirth, C. S. Hartley

A simple graphical technique is presented for rapidly determining the ratio of resolved shear stress on slip systems in single crystals to the applied stress (Schmid factor) for various simple states of stress. Master plots are presented which, when used in conjunction with an appropriate stereo-graphic projection, allow one to determine the Schmid factor for all slip systems in a crystal of given orientation. It is shown that one such plot suffices for simple tension or compression, flex-ural bending, and biaxial compression, while a combination of two different plots must be used for torsion and simple shear. STUDIES of the deformation behavior of single crystals are most valuable when the applied stresses are resolved onto the various slip systems of the crystal. A careful study of the available slip systems and proper selection of the state of stress used and the specimen orientation can then yield quantitative information about the probability of specific dislocation reactions, the strength of barriers to plastic flow, and the mechanisms of yielding and flow in general. It is the purpose of this work to analyze some common states of stress used in deformation studies and to present a simplified way of determining for these cases the ratio of the resolved shear stress on the slip systems to the applied stress (Schmid factor). I) CALCULATIONS We first define two orthogonal coordinate systems: one based on the specimen geometry (a right circular cylinder of radius, R, is considered here) and one based on the particular slip system of interest. The unit vectors and angles specifying the two coordinate systems and the relationship between them are listed in Table I. The relationships among the various angles and unit vectors are shown on a stereographic projection in Fig. 1. The transformation matrix between the (XI X2 X3) and (x&apos;1 x&apos;2 x&apos;3) system is x1 x2 x3 X&apos;1 a11 al2 a13 x&apos;2 a21 a22 a23 x&apos;9 a31 a32 a33 where in this case the direction cosines aij are given by a11 a12 a13 aij = a21 a22 a23 a31 a32 a33 sin 0 cos F cos 0 sin F cos 9 = -cos 0 cos F sin 0 sin F sin 6 [1] 0 -sin F cos (p In terms of Eq. [I.], the stresses sij, referred to the slip system, are determined from oij, referred to the specimen axes, by1,2

Jan 1, 1965

By P. Nyul, S. Caplan, M. F. Lamorte, T. Gonda

Therrnal resistance is measured on GaAs laser diodes in the temperature range 77" to 300°K. These data show that typically the thermal resistance increases fifteen times from 77 to 300°K. The increase is attributed to the decrease in thermal conductivity of Gds. Nornzalized thermal-resistance experimental and calculated from published thermal-conductivity data) values are compared and good agreement shown. In many practical applications, it is desirable to obtain high average power output from laser diodes. Average power output is determined in large measure by the external efficiency and thermal resistance of the diode.&apos;-3 The average input power with which a laser diode may be pumped is dependent on the thermal resistance. Thermal resistance has been studied over the temperature range from 77" to 300°K which is usually of interest in practical applications. This correspondence presents data on thermal resistance for GaAs laser diodes and compares them to values calculated from thermal-conductivity data. Junction heating arises from three distinct components: 1) joule heating, 2) nonradiative recombinations, and 3) reabsorption of radiation. In high-threshold lasers, joule heating predominates.4 Typically, the nonradiative recombination is not negligible and contributes significantly to junction heating. The loss caused by reabsorption may be of the order of the joule-heating loss.&apos; The fabrication and the geometry of laser diodes have been described in other publications.5 The data presented were obtained from diodes which were fabricated from the same region of a wafer to minimize differences. Fig. 1 shows four types of heat sinks employed in this study. In types 1 and 2, the laser cavities are alloyed; pressure contacts are employed in types 3 and 4, and indium is plated on the inside faces to improve the intimacy of contact between the diode and the heat sink. The thermal resistance R is defined by the following relationship: where AT is the junction-temperature rise above the heat-sink temperature for the diode power dissipation Pd. For measurement of thermal resistance, a direct current is passed through the junction, in the range from 1 to 50 pct of threshold. The diode is placed in the laser mode by superimposing a current pulse of low-duty cycle on the direct current.3 Junction-temperature rise is obtained from the time-resolved laser spectrum at the leading edge of the current pulse.&apos; Typically, the laser line shift is 2.5A per OK. The resolution employed in these experiments is 0.5A; therefore, the accuracy in the temperature measurement is 0.2K. Typically, the temperature rise is 2" to 20°K. The power dissipation is the difference between the dc power input, and the radiation power exiting the diode. The measurement accuracy is less than 1

Jan 1, 1968