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By G. W. Thompson

G. W. THOMPSON,* Brooklyn, N. Y.-This subject has two aspects, neither of-which can be ignored: these are the economic aspect and the technical aspect. Under ordinary conditions, economic law will take care of the conservation of tin. Under present conditions it seems desirable that economic law should still be permitted to operate as far as is practicable. It is true that this law operates rather slowly, and that under war conditions it cannot be depended upon to give our Government the supply it immediately needs. There should, therefore, be a designation of the essential industries by the Government, and tin should be supplied for those needs, letting the non-essential industries take what is left. At the same time, information ought to be forthcoming as to how tin can be conserved; with tin selling at 80 c. per pound, more or less, it is to the interest of every consumer to get the best information he can as to how he can save tin. Economic law would properly punish those who do not study their own interests in this way. I deprecate, therefore, any centralized socialistic effort toward avoiding penalties through failure to observe and obey economic law. In saying this, however, I am not unmindful that it is desirable for our Government to take such control of the tin situation as may be necessary to the prosecution of the war, practically regardless of the effect of such action upon individual industry, which should studiously seek to adjust itself to the new condition. The Government should promulgate such technical information as it can collect, showing how tin can be conserved, and should urge upon every consumer the exercise of his common sense in self-protection. There is very little danger of any one attempting to hoard or corner tin, in the present state of the market. Most consumers will be glad to live from hand to mouth, covering their sales by purchases, or vice versa. As to the technical aspects: Tin and tin alloys are used to give certain practical and also certain artistic results. The practical factors are those involved in proper adhesion, continuity of surface, protection, the right degree of hardness, proper working qualities, etc. The artistic results are those that appeal purely to the eye or to our cultivated sense of what is desirable. There is no doubt that a great deal of tin could be saved if there were not a demand for certain pleasing effects. How can this saving be brought about? The users of tin hearing alloys are not the only ones involved in this question. The manufacturers of solder and babbitt have for a long time sought to give certain appearances to their fabricated bars and ingots in order to make them more salable. The user of a solder is very apt to judge of its working qualities by the appearance of the bar; the same thing is true with regard to babbitts. Under ordinary conditions, manufacturers are justified in trying to

Jan 12, 1918

By K. L. Young

The general energy equation, including change in kinetic energy, was solved by numerical integration and used to evaluate simplifying assumptions and application practices over a wide range of conditions. When extreme conditions were encountered, sizable errors were caused by large integration intervals, application of Simpson's rule and neglecting change in kinetic energy. A maximum error of only 1.31 percent was caused by assuming temperature and compressibility constants at their average value. It was discovered that a discontinuity can develop in the integral for the injection case. This discontinuity indicates a point of zero pressure change and is an inflection point in the pressure traverse. INTRODUCTION When a pressure in a gas well is to be calculated, one of the first decisions is to select a method of calculation. In many instances, this selection becomes a problem because the literature, at best, provides an evaluation of any method for only a limited range of conditions. Once a method has been selected, a question often arises as to the size of the calculation interval which should be used. The question regarding calculation interval arises because an analytic solution is not obtainable and approximate solutions must be used. This paper presents an evaluation of major assumptions and application practices of probably the two most widely used methods for calculating steady-state single-phase gas well pressures. The two methods are Cullender and Smith1-" (numerical integration), and average temperature and com-pressibility. The Cullender-Smith method assumes that change in kinetic energy is negligible and is normally applied in two steps with a Simpson's rule correcton. The average temperature and compressibility method, in addition to neglecting kinetic energy change, asumes that temperature and compressibility are constant at their average values. This method is normally applied for wellhead shut-in pressures of less than 2,000 psi, and in one step. Computer programs were written to compute bottom-hole pressure with and without the assumptions, using various approaches. Values of input parameters investigated are shown in Table 1. Flow rate was limited to a maximum of 5,000 and 10,000 Mcf/D for tubing sizes of 1.610 and 1.995 in. ID, respectively. Flow rate was also limited to 10,000 Mcf/D for a tubing size of 2.441 in. ID when wellhead flowing pressure was 100 psia. These limitations were imposed on flow rate so as not to exceed sonic velocity. The z factor routine available necessitated limiting bottom-hole temperature to 240F and wellhead pressure to 3,000 psia. Pressures were compared on the basis of percent deviation from the trapezoidal integration of Eq. l or 2 at 100-ft intervals. A preliminary investigation indicated that a 1,000-ft interval solution would differ from a 50-ft interval solution by less than 0.25 percent; therefore, the 100-ft interval was chosen for a base. For the purpose of comparison, deviations less than 1 percent were considered insignificant. where 111.1 qeC2/d p = kinetic energy term. Eqs. 1 and 2 can be evaluated numerically at specific depths using the trapezoidal rule as shown by Cullender and Smith. If change in kinetic energy is neglected and temperature and compressibility are assumed constant at their average values, Eq. 1 can be integrated to give the average temperature and compressibility equation,

By Jed H. Mosgrove, Paul M. Budzak

INTRODUCTION Good mine ventilation starts at the drawing board. A coal mine will either be adequately ventilated or not, depending upon how good a job was done in the planning. Poor planning will not only result in inadequate ventilation, but may also lead to low productivity. A mine foreman cannot achieve high productivity while trying to maintain 4.2 m3/s (9000 cfm) of air at his last open crosscut. PLANNING What will provide adequate ventilation? A rule of thumb in laying out a coal mine is to provide enough entries in the main entry system to keep air velocities at 2.5 m/s (500 fpm) or less; at this rate of flow, good ventilation may be expected. The intake and return air courses of a coal mine do not improve with age. Entries will deteriorate, ribs will slough, roof will fall, and clutter will accumulate. The resistance offered to air travel will continue to increase with each passing year. Providing enough airways for the present and the future cannot be emphasized too strongly. In order to better appreciate good planning and bring it into better focus, assume a hypothetical mine has the following features: entries 1.5 m (5 ft) high and 6.1 m (20 ft) wide; one intake and one return entry; the mine development is 4.8 km (3 miles) from the intake and upcast ventilation shafts; and the ventilation requirement is 141.6 m3/s (300,000 cfm). Air travel of 4.8 km (3 miles) into a mine and the same distance back to the fan is not unusual. However, to do so with one intake and one

Jan 1, 1981

By Jed H. Mosgrove, Paul M. Budzak

INTRODUCTION Good mine ventilation starts at the drawing board. A coal mine will either be adequately ventilated or not, depending upon how good a job was done in the planning. Poor planning will not only result in inadequate ventilation, but may also lead to low productivity. A mine foreman cannot achieve high productivity while trying to maintain 4.2 m3/s (9000 cfm) of air at his last open crosscut. PLANNING What will provide adequate ventilation? A rule of thumb in laying out a coal mine is to provide enough entries in the main entry system to keep air velocities at 2.5 m/s (500 fpm) or less; at this rate of flow, good ventilation may be expected. The intake and return air courses of a coal mine do not improve with age. Entries will deteriorate, ribs will slough, roof will fall, and clutter will accumulate. The resistance offered to air travel will continue to increase with each passing year. Providing enough airways for the present and the future cannot be emphasized too strongly. In order to better appreciate good planning and bring it into better focus, assume a hypothetical mine has the following features: entries 1.5 m (5 ft) high and 6.1 m (20 ft) wide; one intake and one return entry; the mine development is 4.8 km (3 miles) from the intake and upcast ventilation shafts; and the ventilation requirement is 141.6 m3/s (300,000 cfm) . Air travel of 4.8 km (3 miles) into a mine and the same distance back to the fan is not unusual. However, to do so with one intake and one

Jan 1, 1981

By Oliver S. North

Reserves of perlite rock in the western section of the United States are immense. A geological report prepared for the Union Pacific RR showed proved tonnage of over 400 million tons in southern Nevada alone. Yet the perlite industry is not large, and most of the mines are relatively small. There are indications that perlite processing, thus far largely in the hands of relatively small com¬panies, is becoming concentrated in a few large companies. For economic reasons the bulk of perlite sold in crude form has come from three or four pro¬ducers, although many individual operators have sold small quantities to limited lists of purchasers. Perlite is one of ultra-lightweight aggregates. Owing to its outstanding qualities as an aggregate, variety of uses, and ultimate economy of use, it has enjoyed considerable popularity in this country in recent years. Marked foreign interest also has been evinced in this material.

Jan 2, 1955

By Babbitts and Solder

G. w. Thompson,* Brooklyn, N. Y.—This subject has two aspects, neither of which can be ignored: these are the economic aspect and the technical aspect. Under ordinary conditions, economic law will take care of the conservation of tin. Under present conditions it seems desirable that economic law should still be permitted to operate as far as is practicable. It is true that this law operates rather slowly, and that under war conditions it cannot be depended upon to give our Government the supply it immediately needs. There should, therefore, be a designation of the essential industries by the Government, and tin should be supplied for those needs, letting the non-essential industries take what is left. At the same time, information ought to be forthcoming as to how tin can be conserved; with tin selling at 80 c. per pound, more or less, it is to the interest of every consumer to get the best information he can as to how he can save tin. Economic law would properly punish those who do not study their own interests in this way. I deprecate, therefore, any centralized socialistic effort toward avoiding penalties through failure to observe and obey economic law. In saying this, however, I am not unmindful that it is desirable for our Government to take such control of the tin situation as may be necessary to the prosecution of the war, practically regardless of the effect of such action upon individual industry, which should studiously seek to adjust itself to the new condition. The Government should promulgate such technical information as it can collect, showing how .tin can be conserved, and should urge upon every consumer the exercise of his common sense in self-protection. There is very little danger of any one attempting to hoard or corner tin, in the present state of the market. Most consumers will be glad to live from hand to mouth, covering their sales by purchases, or vice versa. As to the technical aspects: Tin and tin alloys are used to give certain practical and also certain artistic results. The practical factors are those involved in proper adhesion, continuity of surface, protection, the right degree of hardness, proper working qualities, etc. The artistic results are those that appeal purely to the eye or to our cultivated sense of what is desirable. There is no doubt that a great deal of tin could be saved if there were not a demand for certain pleasing effects. Row can this saving be brought about? The users of tin bearing alloys are not the only ones involved in this question. The manufacturers of solder and babbitt have for a long time sought to give certain appearances to their fabricated bars and ingots in order to make them more salable. The user of a solder is very apt to judge of its working qualities by the appearance of the bar; the same thing is true with regard to babbitts. Under ordinary conditions, manufacturers are justified in trying to

Jan 1, 1919

By David Segura

THE Huitzuco mercury deposits have been described by C. W. Vaupell,1 therefore the general information such as location, history, production, etc. will not be repeated here. The present owners acquired the property when the deep levels of the Cruz mine were flooded, and it was believed that a relatively large tonnage of low-grade ore and fills would be encountered in the flooded section, because of the early mining methods used in the Cruz mine. The mine was partly unwatered and about 187,000 tons of fills was pit-sampled in the San Cayetano stope. These fills average 2.4 kg. Hg per ton. During the period of maximum production at Huitzuco, the ores were sorted in the mine and cobbed into 3-in. pieces, which were treated in Scott furnaces. Large pieces of raw ore are found in the sinter, and it is indicated that both recovery and costs were not very satisfactory. Recent practice has been to treat the ores in closed retorts made of 10-in. pipes. Because of the presence of combined antimony in the mineral livingstonite (HgS.2Sb2S3) and considerable native sulphur, little or no metallic mercury is recovered in the condensers and a second retorting of the condenser muds is necessary. This makes the treatment of the ores expensive and tends to low recoveries of mercury. The principal costs of retorting are, in their relative order: (1) fuel, (2) low recovery of mercury, (3) replacement of retort pipes, (4) labor. METHODS TESTED Three series of experiments have been made, with the object of reducing the costs of recovering mercury and antimony from the ores, by: (1) wet metallurgical methods, (2) bulk roasting in large mechanical furnaces, (3) concentration to reduce the bulk to be followed by roasting in some mechanical furnace. The wet metallurgical methods did not appear feasible because of the high cost of the alkaline solvents used, and investigations of these were not carried beyond the preliminary stage.

Jan 1, 1938

By P. D. Merica

Certain compositions of the light, i:e., aluminum-rich, alloys of aluminum with magnesium and copper have become quite well known within the past ten years under the name of duralumin. These alloys are used for rolling and forging and can be so treated as to develop quite remarkable mechanical properties. Thus a properly heat-treated alloy containing about 4 per cent. of copper and about 0.5 per cent. of magnesium, rolled into sheet or rod, will have a tensile strength of approximately 55,000 lb. per sq. in. (3866 kg. per sq. cm.) with an elongation in 2 in. (50.8 mm:) of about 15 per cent. This alloy, duralumin, was developed by Wilm,1 and its properties more fully described by Cohn.2 The authors considered it worth while to investigate the mechanical possibilities of the light alloys of two somewhat analogous ternary series; namely, of aluminum-magnesium-nickel and of aluminum-magnesium-manganese, to study the effect of variation of composition upon mechanical properties within the alumimum-rich group of the aluminum-magne-

Jan 7, 1919

By A. S. Yue

The movphology of the Mg-32 wt pct Al eutectic has been studied as a function of freezing- rate and temperature gradient. At slow freezing rates a lamellar eutectic was formed; whereas, a rod-like eutectic was generated at fast rates. The inter-lamellar spacing increased as the freezing rate decreased in aggreement with theoretical predictions. Lamellar faults, morphologically similar to edge dislocation models in crystals, were responsible for the subgrain structures in the eutectic mixture. A linear increase in fault density with freezing rate was observed. Fault concentl-ations of the order of 10 per sq cm for a range of freezing rates from 0.6 to -3.0 x 10 cm per sec were estimated. The transformation from lamella?, to rod-like morphologies was determined experimentally to be dependent on the freezing rate and independent of the temperature gradient. Moreover, the number of rods formed per- unit cross-sectional area increased exponentiallv with increasing freezing rote. BRADY' and portevin2 classified eutectic structures into lamellar, rod-like, and globular according to the morphology of the solid phases present. Although this classification is quite descriptive, very little has been reported on the details of the mechanism by which the eutectic structures are formed. Recent work by Winegard, Majka, Thall, and chalmers3 and by chalmers4 on lamellar eutectic solidification suggest that the maximum thickness of the lamellae decreases with increasing rate of solidification due to inadequate time for lateral diffusion. scheilS and Tiller' have shown theoretically that the lamellar widths indeed depend on the solidification rate. However, there has been no experimental evidence to support the theory. Chilten and winegard7 have studied the interface morphology of a eutectic alloy of zone-refined lead and tin. They found that the lamellar width decreased as the freezing rate increased in agreement with the theoretical predictions of scheils and Tiller.' More recently, Kraft and Albright' have investigated the microstructures of the A1-CuA12 eutectic as a function of growth variables. They observed lamellar faults present in the lamellar eutectic, similar to edge dislocation models in crystals. Furthermore, Kraft and Albright reported that they could not designate which extra lamellar was responsible for the formation of a lamellar fault even under electron microscopic magnification. In this paper, the morphology of the Mg-A1 eutectic structure is described. The effects of freez- ing rate on the interlamellar spacing and on the lamellar fault density are presented in detail. The transformation from lamellar to rod-like eutectics is discussed in terms of the freezing rate and the temperature gradient. EXPERIMENTAL PROCEDURE The experimental details of alloy preparation, the decanting mechanism and the determinations of the freezing rate and the temperature gradient have been reported elsewhere. Measurements of plate-edge angles were made with a microscope. The true angles used to determine the interlamellar spacings were determined by a two surface analysis technique.'' Since the decanted interface structure does not represent the true eutectic morphology on the solid,g all measurements were made from an area in the solidified bar behind the interface. Measurements of the apparent interlamellar spacings between the two phases of the eutectic were made on a photographic negative by means of a calibrated magnifier. Each value listed in Table I represents the average of thirty measurements on one negative. In general, these measurements are approximately equal with an error of less than pct. The average rod diameter for each specimen was also measured on a magnified photomicrograph. Each value of the diameter represents the average of fifty measurements. RESULTS AND DISCUSSION The experimental observations and their discussion to be presented here are restricted to the morphology of the eutectic structure and to the effects of the freezing rate and the temperature gradient on the solidification of eutectics. INTERLAMELLAR SPACING It has been shown previouslyg that the micro-structure of the decanted interface and the longitudinal section of the Mg-A1 eutectic is characterized by the presence of both lamellar and rod-like morphologies. The lamellae become more regular as the freezing rate is decreased. A three-dimensional photomicrograph representing a perfect lamellar morphology is illustrated in Fig. 1. The lamellae of the top and longitudinal sections of the specimen are regularly spaced while those in the transverse section are not quite straight and parallel. Their parallelism is slightly distorted because fault lines producing a discontinuity are present. A method for calculating the interlamellar spacings A, is described in Appendix 1. The true

Jan 1, 1962

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

The hardness dependence on grain size for polycrys-talline cartridge brass and a leaded brass has been measured by Brine11 and Rockwell B testing. In each case, the hardness, H, depends on the average grain diameter, 1, according to: H =Ho + kHl-1/2 where Ho and kH are experimental constants. Diamond pyramid hardness values have also been measured as a function of the indentation size and grain size to give additional information on the nature of the hardness test and the dependence of hardness on micro-structure. The hardness of polycrystalline brass depends on its grain size. Bassett and Davis' demonstrated this as early as 1919 by making Brinell hardness measurements on cartridge brass. Since then, the hardness of this type of material has been measured as a function of grain size by making Rockwell,2'3 Vickers,4 and Brinell5 tests. he hardness dependence on grain size has also been measured for other materials. Angus and summers6 investigated the grain size dependence of the Brinell hardness of polycrystalline copper and a Cu-4.5 pct Sn bronze. In other studies, nickel,? Armco iron,Big an Fe-0.07 pct C alloy,I0 and an 0.39 pct C-12.45 pct Cr stainless steel" have been investigated. In some of the preceding cases, the hardness results have been analyzed to show that the hardness varies with the average grain diameter, 1, according to an l-l\4, l-1/4 or I-2 dependence,11-13 The studies of the influence of grain size on hardness have not been based on any theoretical model. This may be because the hardness of a material is itself a complicated property. However, attempts have been made to correlate, experimentally and theoretically, the hardness of a material with its unidirectional stress-strain behavior.14-l6 On this basis, Hall" proposed that the polycrystal hardness dependence on grain size might follow directly from the Hall-Petch18,19 relation for the grain size dependence of the yield stress. Thus, the hardness-grain size relation was given as: H = Ho + kHl-1/2 [1] where Ho and kH were taken as experimental constants. The relation was applied to the measurements on brass,' copper,6 bronze,= and Armco iron.' More recently, this relation was shown by Armstrong and jindal20 to adequately describe the measurements on cartridge brass made by Bassett and Davis' and Babyak and Rhines.5 In this case, the relationship was taken a step further by independently relating the values of Ho and kH to the values of oyand ky, previously reported by Armstrong, Codd, Douthwaite, and petch21 from measurements of the yield stress dependence on grain size for this type of material. In the present investigation, new Brinell and Rockwell B hardness measurements have been made as a function of grain size for a cartridge brass and a leaded brass. In addition, diamond pyramid hardness values were measured as a function of the indentation size. All these results are applied to a further analysis of the hardness dependence on grain size. MATERIALS AND EXPERIMENTAL METHODS Cartridge brass and a leaded brass were selected for this investigation for two main reasons: it was anticipated 1) that these materials could be cold-worked and recrystallized to a wide range in grain size and 2) that the results to be obtained on these typical industrial materials could be usefully compared with previous investigations. The chemical analyses of the actual materials which were employed are given in Table I. The as-received 1/2- and 3/4-in.-thick plates were given various reductions in thickness by cold rolling. The rolled material was heat-treated at various temperatures between 330" and 850°C for differing time periods from 5 min to 9 hr to achieve a variation in the average grain diameter between 0.0339 and 0.000543 cm.22 During heat treatment, the brass was protected from zinc loss by packing it in chips or foils of the same composition material. Reasonably equi-axed grain structures were obtained in each case. The metallurgical grain sizes of the specimens were determined from measurements of the average linear intercept on a random line. Annealing twin interfaces were not counted along with grain boundaries. The

Jan 1, 1970

By Y. C. Liu, G. A. Alers, S. S. Choi

In searching for an explanation for the obvious difference between the copper and the brass type of rolling texture, it is common practice to look for correlations with other properties. For example, there are obvious differences between copper and brass in the critical resolved shear stress, the extent of easy glide or stage I, and the amount of "overshooting)) during single-crystal deformation.' Also, the slip lines in copper are fine and wavy while those in brass are more coarse and better described as slip bands.' Since neutron irradiation is known to make these properties in copper change in such a way as to resemble brass,2'3 an excellent test for these possible correlations is to see whether neutron irradiation produces a corresponding change in the texture, i.e., whether irradiated copper develops a brass-type rolling texture. The experiment would also answer the more important question of whether or not lattice defects like those introduced by neutron irradiation have any effect on the development of rolling textures in fcc metals. Experimental studies using conventional X-ray pole figure techniques are inconvenient because the X-ray detection system must be modified to discriminate against the radioactivity which is induced in the specimen by the necessarily long exposure to neutrons. Nevertheless, some texture determinations have been

Jan 1, 1968

By R. F. Marande, R. S. Busk

THE magnesium-aluminum-zinc system contains most of the magnesium-base alloys used commercially, although in practice the ternary, alloys are usually modified by the addition of a small amount [ ] of manganese. Table I contains the magnesium-base commercial sand-casting alloys most widely used throughout the world, together with their typical tensile properties. These alloys may be regarded as a single family of varying aluminum and zinc content, disregarding manganese. A concise means of presenting data for them is by means of iso-property lines plotted on triangular coordinate paper. Constitutional, casting, and mechanical properties are presented in this manner. The data are given with the object of providing reference material, and to aid in the choice of a commercial alloy to serve a particular purpose. Detailed, critical analysis of the commercial alloys has been covered in a previous paper.1 [ ] All alloys reported were made in 15-lb. melts, superheated to 1700°F. for 15 min., cooled to 1400°F. and poured in sand molds. The test bars produced had a ½ in. reduced section and were about 6 in. long. Except for the data represented in Figs. I and 3 the alloys used were of commercial purity and contained 0.1 to 0.2 per cent Mn in addition to the aluminum and zinc. CONSTITUTIONAL DATA Under the heading of Constitutional Data are included all known data on con-

Jan 1, 1946

By Eugene F. Broome, Hugh A. Walls

Self-diffusion coefficients were measured using the modified shear cell technique over a temperature range from 31" to 401°C. These data agree within experimental error with those of Petit and Nachtrieb, which were determined from 30°to 99°C. LIQUID self-diffusion behavior in gallium has been studied previously by Petit and Nachtrieb&apos; using their shear cell technique.&apos; Data were obtained in that investigation over temperatures ranging from 30.6 to 98.7"C, and these comprise the only published data for liquid gallium. In the work reported here, diffusion data were obtained from 30.8 to 400-9.Cusing a shear cell technique different from that of Petit and Nachtrieb. This modified shear cell method and the mathematical analysis for diffusion in a finite capillary geometry have been previously described.3 PROCEDURES AND ANALYSIS The segments of the modified shear cell were fabricated from types 304 and 316 stainless steel. No corrosion problems with gallium were encountered until attempts were made to obtain data at 500°C. The diffusion capillary diameter was 1.59 mm and capillary lengths for the four sets of cell segments used are given in Table I. Gallium with a purity of 99.999 wt pct was obtained from Var-Lac-Oid Chemical Co. The gallium-72 isotope used to measure the self-diffusion coefficients was produced from this inert gallium by irradiation in The University of Texas Reactor Facility, and was checked for radiochemical purity. Following a period of isothermal diffusion controlled to within 0.0l°C, the contents of each capillary section were emptied into a counting planchet, weighed to assure that gallium had not been lost during handling, and the radioactivity determined by windowless gas-flow proportional counting. Great care was taken to obtain a reproducible counting geometry for the samples. Diffusion coefficients were obtained from the solu-tion3 to the diffusion equation for the modified shear cell problem, viz.: in which B2= C2(t)Li + L2)- CSL1-COL2 P* (C0-Cs)(L1+ L2) and L, and L2 are the capillary segment lengths. The step-function initial concentration distribution was expressed as C(x,O) = C, for -L1 = x < 0 and C(x,O) = Co for 0 < x 5 L?. Also, C2(0 = ^- fcix,t)dx Computation of the diffusion coefficients was done numerically as described in our earlier paper.3 Because of the short half-life of gallium-72 (14.1 hr) and the relatively long counting times taken to give a precise value of the diffusion coefficient, it was necessary to account for the radioactive decay both during the actual count and between counts for different samples (before and after diffusion). Thus, the average count rate Cz(t) for a second sample counted from times t3 to t4 can be made consistent with the average rate Co determined for a first sample counted from times t1 to t2 by the relationship: N(e-»i-e-M») , , C2(f) " &-«,)(*-«.-«-»«) [21 in which N is the total number of counts taken and is the decay constant. RESULTS AND DISCUSSION The self-diffusion coefficients determined in this investigation are listed in Table 11. Values of the parameters used to calculate the coefficients are given in the table, as are values of the random counting error &apos; for the concentration ratio pz and the corresponding expected random variation AD in the diffusion coefficient. The quantity is calculable from: &apos;Cg — C2(t)} , „ ,2 .-„¦, M co-csl(ACs)J W in which AC is the error in the individual count rate. The mean values of the diffusion coefficients within each temperature group are also given in Table 11. Such temperature groupings of data are reasonable because of the small variations between the temperature~.~

Jan 1, 1970

By A. R. Marder

Fe-C alloys have been investigated at temperatures below the eutectoid transformation to determine whether the superplasticity phenomenon exists for these materials. As a result of void formation at the carbide-ferrite interfaces that led to void coalescence and premature failure, the large elongations characteristic of superplustic matem&apos;als did not develop. It had generally been assumed that the conditions for obtaining superplasticity include: a) T/Tm =0.5, b) a strain rate sensitivity of the flow stress that results in m s 0.3, and c) fine-grain structure. Although superphsticity was not observed in the alloys studied, stress-strain rate measurements indicated that an additional condition for the existence of superplasticitv in Fe-C allovs is that the fine-wain structure be equiaxed in agveement with previous results on Zn-Al al1oys. MANY investigations have been made on the hot deformation of iron-base alloys in order to obtain an understanding of the variables that may affect high-temperature strength and ductility.&apos;-6 However, only a few papers have studied the effect of carbon content on the high-temperature properties of iron-base alloys. Peretayat&apos;ko and Zaikov and Wallquist and carlen4 studied the effect of carbon on strength in the y temperature range, and recently Robbins, Shepard, and sherby6 studied the influence of carbon on the torsional behavior of iron in the a and cu-plus-cemen-tite temperature ranges. In addition to studying hot deformation behavior, current investigations have also been concerned with the phenomenon of superplasticity in nonferrous materials, and, in particular, ferrous alloys. Superplasticity is identified by the increased resistance to necking that results from a high strain rate sensitivity of the flow tress. ackofenr proposed the strain rate hardening coefficient m = d log a/d log e as a measure of strain rate sensitivity and showed that m must be "-0.3 for superplasticity to exist. In addition, a fine-grained microstrcture,- a eutectic or eutectoid ornosition,- and a test temperature in the creep range (TITm > 0.5)" have also been advanced as necessary conditions for superplasticity. Although many workers""23 investigated superplasticity associated with the allotropic transformation in iron and the eutectoid transformation in steel, and recently chadler " and orrison have studied superplasticity in the a - y phase region, no one has measured the superplastic effect in the a + cementite temperature region. Therefore, the purpose of this investiga- tion was to determine the effect of carbon content, microstructure, and test temperature on the strain rate sensitivity of Fe-C alloys at temperatures below the eulectoid transformation. EXPERIMENTAL PROCEDURE Fe-C alloys with carbon contents of 0.2, 0.4, 0.6, 0.8, and 1.0 wt pct were prepared as 100 lb 54 in. ingots by vacuum-induction melting. The compositions are listed in Table I. All the ingots were soaked at 2100°F for 2 hr prior to rolling to 4 in. plate. The minimum finishing temperature for rolling was 1800°F. The plate was sheared in 10-in. lengths and tensile bars were rough-machined prior to heat treatment. The samples were heat-treated according to the schedule in Table I1 in order to obtain a fine-grained microstructure consisting of finely dispersed carbides in a matrix of small ferrite grains. All the tensile specimens were heat-treated prior to finish machining, and tensile testing was accomplished on an Instron machine in conjunction with a clamshell furnace. The buttonhead specimen configuration is shown in Fig. 1. An inert atmosphere, helium, was maintained during testing with a steel tube shroud in a bell jar arrangement. Predried helium was continually purged over the buttonhead specimens and exhausted through ceramic fiber at the bottom of the tube to keep oxidation at a minimum and to prevent

Jan 1, 1970

By H. I. Aaronson, S. Toney

When the component grains of .ferritic hicrystals of an Fe-1.55 pct Si alloy are disoriented through an angle "6 " about a conzmon [ll0] axis, the tendency for preferential growth of austenite crystals along the grain boundary during transformation at elevated temperatures is small when 11 deg, but increases rapidly at larger angles. This type of orientation-dependence indicates that grain boundary diffdsion promotes preferential growth along large-angle boundaries. Morphological differences between austenite crystals formed at small-angle [1101 and [loo] boundaries suggest that precipitate morphology can be dependent on the dislocation structure of the boundary. ThE morphology of precipitate crystals nucleated at a grain boundary can be significantly affected by the structure of the boundary.&apos; The limited amount of experimental evidence available in the literature indicates that the morphological effects of boundaries made up of arrays of dislocations, such as subbound-aries and small-angle grain boundaries, are different from those of boundaries having essentially disordered structures, i.e., large-angle grain boundaries. On the basis of indirect evidence, it has been concluded that large-angle grain boundaries give rise to the formation of grain boundary allotriomorphs (crystals which nucleate at grain boundaries, and grow preferentially and more or less smoothly along them)2 in the proeutectoid ferrite and the proeutec-toid cementite reactions in plain-carbon steels, and apparently also in many non-ferrous alloys.&apos; At small-angle grain boundaries in a plain carbon steel, on the other hand, ferrite crystals were found to take the form of primary side plates. Similarly, Guinie, loys. Primary sideplates formed at a subboundary with a constant orientation tend to be parallel to only one, or occasionally two matrix habit planes, and a marked change in the orientation of the boundary is accompanied by a change in the habit plane. Previous studies on the morphological effects of grain boundary structure were performed on poly-crystalline aggregates. Information on the disorien-tation of the pairs of grains forming the boundaries at which the various morphologies appeared in these specimens was largely either qualitative or semi-quantitative. Precipitate morphologies accordingly could not be accurately and systematically correlated with grain boundary structure, and thus theories which have been proposed for the various morphological effects could not be satisfactorily tested. This investigation was undertaken in an attempt to remedy these deficiencies by studying the morphological effects of grain boundary structure with a method in which the boundaries are formed by matrix grains whose disorientations are known and controlled with reasonable accuracy. EXPERIMENTAL PROCEDURE The crystallographic requirements of this study were fulfilled by means of oriented bicrystals of silicon-iron. Disorientation of the component ferrite crystals was carried out about common, major crystallographic axes through angles ranging from 1/2 to 44 deg. The silicon content was low enough SO that the bicrystals could be partially transformed to austenite by heating to elevated temperatures. The silicon-iron used had the following initial composition: 1.55 pct Si, 0.04 pct C, 0.0031 pct N, 0.17 pct Mn, 0.020 pct S, and 0.002 pct P. The alloy was obtained in the form of 0.036-in. sheet. The procedures employed to prepare seed crystals in strips of this sheet, to reorient the seeds, and to grow them into bicrystals are essentially those described by Dunn and Nonken and aynes." The characteristics of the bicrystals are given in Table I. The "bicrystal type" indicates the crystallographic plane parallel to the broad faces of the strip in both grains and the crystallographic direction which was parallel to the long edges of the strip in both grains prior to disorientation. The angular disorientation of the grains, 8, which was performed about the direction normal to the plane of the broad faces of the strip, was measured between the [ 00l] directions. Orientation of the grain boundary, , was taken as the angle between the plane of the grain boundary and a plane containing the axis of disorienta-

Jan 1, 1962

RICHARD R. HICE, Beaver, Pa. (written discussion *).-The matter of selective. mining is probably of more importance in Pennsylvania than is washing, and perhaps washing would not be necessary at some plants where now practised if proper care were taken in the mining. I have in mind one plant where the sulfur content of the entire bed is more than 2 per cent., yet if the bottom 12 in. and the top 12 or 15 in. are separated, the main portion of the bed, some 6 ft., will not contain more than 1.25 per cent. sulfur; and yet in this mine the entire bed .is worked together, the resultant coke, as a matter of course, not being fit for use in iron smelting. Not sufficient notice is taken of the "Double-thick" Upper Freeport coal in Allegheny and Butler counties. This is a very valuable deposit. Other areas where there is a, top member of the Upper Freeport coal are also known, which add to the value of this bed as a coking coal. This double Freeport coal is especially suited for byproduct coking. I cannot be as optimistic as regards the "low-sulfur" Pittsburgh coal in Greene county as the authors. If the term low-sulfur is to he confined to coal with not exceeding 1.25 per cent. sulfur, there is probably little

Jan 10, 1919

By W. A. Backofen

THE texture of cold-drawn copper wire has been described most often as a composite of [Ill] and [loo] directions aligned parallel to the wire axis.&apos; Hibbard2 suggested that sufficient reduction will result in a single [111] texture and reported that after a reduction in area of 96.4 pct, such a texture was nearly obtained. Other work3-4 has shown that recrystallization within the temperature range of 300" to 1000°C does not alter the texture established by cold-drawing. In general, however, the texture of cold-drawn end annealed copper wire is not as well defined as the deformation texture. In the work reported here it was found that the double-fiber, deformation texture persists even after a reduction in area of 97.3 pct, and after recrystallization the texture is practically all [l00]. The wire used in this work was prepared in the following way. A 3/4-in. diam length of annealed OFHC copper was cold-drawn to a diameter of 1/8 in., a reduction in area of 97.3 pct. Each step in the drawing operation brought about a reduction in area that varied from 6 to 10 pct. Short lengths of this wire were etched in a 50-pct solution of nitric acid to a diameter of approximately 0.01 in. The texture of the cold-drawn specimens was determined with the conventional X-ray transmission diffraction technique using filtered copper radiation. The X-ray beam was directed perpendicularly to the wire axis and the diffraction rings were recorded on a flat film. Fig. la is typical of the photo-grams obtained and reveals that the cold-drawing schedule described above introduces a distinct, double-fiber texture with both [Ill] and [l00] directions aligned parallel to the wire axis. Supplementary measurements with a Geiger-Mueller counter of the intensity of reflection from (111) planes showed that the [111] component was much more predominant than the [l00]. Additional specimens of the drawn and etched wire were annealed in a salt bath at temperatures from 300" to 450°C. The texture of the recrystal-lized wire was studied with the same technique applied to the cold-drawn wire. Fig. lb is an X-ray photogram of a specimen annealed for 1 ½ at 300°C and is typical of those obtained from re-crystallized specimens. The photogram shows that recrystallization of this material is accompanied by the almost complete elimination of the [111] component of the deformation texture. A pronounced strengthening of the [100] component was evident from intensity measurements with a Geiger-Mueller counter. In summary, the deformation texture of cold-drawn OFHC copper wire, after a reduction in area of 97.3 pct, may still be described as a composite of [Ill] and [loo] directions aligned parallel to the wire axis. Recrystallization of such wire at temperatures from 300" to 450°C introduces a texture which consists almost entirely of a [loo] alignment. Acknowledgment Sidney Rolle of the Scomet Engineering Co. generously supplied the OFHC copper used in this work. References &apos;C. S. Barrett: Structure of Metals. p. 382 (1943) New York. McGraw-Hill Book Co. 2 W. R. Hibbard, Jr.: Trans. AIME (1949) -185, p. 598; Journal of Metals (Sept. 1949) TN 19E. 3 P. 422 of ref. 1. &apos;E. Schmid and G. Wasserman: Ztsch. Physik (1926) 40, p. 451.

Jan 1, 1952

R. Culver, Colorado School of Mines I appreciated Professor Starfield&apos;s presentation, because his approach to blasting research reflects my feelings that a lot of practical work can be done toward making short-term advances that can be used by the blasting industry. We should not forget this in our attempt to obtain solutions to such idealized problems as a mathematical theory of cratering. The blasting industry would like very much to have the efficiency of the basic blasting processes improved from the range of 20 percent up into the range of 50 percent. And many of these problems should be approached and attacked in the laboratories. Blasting research at the Colorado School of Mines for the past two or three years has been aimed at improving the pre-splitting process, because we feel it will become increasingly important to the construction industry. We have attempted to develop a unified theory of how pre-splitting works; and the more we investigate, the more variables we find. We may never obtain an entirely satisfactory theory of pre-splitting, but we have developed a tentative explanation. One of our first conclusions was that we are making very poor use of the explosive in current pre-splitting practice. But to appreciate this we must first consider how pre-splitting works. A series of parallel boreholes are loaded with low density charges which are detonated simultaneously. (Delayed detonation has been tried, but most people agree that simultaneous detonation is best, and this is what we found.) Stress waves traveling out from each borehole interact with the stress waves from the two adjacent boreholes influencing the direction of the fractures that occur and propagate from each borehole. It has been suggested in private communication with Mr. Wilbur Duvall of the U. S. Bureau of Mines, that the fracture between the adjacent boreholes is a result of the net particle velocities above and below the pre-split plane which tend to set up a net tension along the pre-split plane. Regardless of what the exact mechanism is that causes the fracture to propagate from one borehole to the next, we know that it, in general, relies upon an interaction of the stress waves. However, if there is a large compressive stress normal to the pre-split place, it can practically eliminate any possibility of obtaining the desired pre-split by preventing the desired fractures from getting started. Aside from obtaining a fracture along the pre-split plane, the other major aim of pre-splitting is to minimize the fracturing that occurs around the borehole. We do this by drilling a large hole and putting a small charge in the middle of it. This results in a badly decoupled charge which is very inefficient; most of the energy is used in compressing the air in the borehole. Out of discussions about how pre-splitting works, we decided that perhaps there was some other explosive geometry which would make more efficient use of the explosive, and result in optimizing the pre-splitting process. Starting in the laboratory with small plexiglas blocks, we found that a fully-coupled explosive charge of 5 grains/in. of PETN gave random fractures extending about I in. from the borehole. A 1/4 in. plexiglas plug fit into a 1/4 in. hole in a similar block with the same fully coupled charge gave random fractures extending out only about 1/2 in. due to the energy absorbed in fracturing the plug. If the plug was split longitudinally but otherwise using the same geometry, it was possible to drive a fracture about 2 in. in either direction from the borehole, with an even smaller amount of fracturing around the borehole.

Jan 1, 1967

By P. E. Busby, C. Wells

FURTHER study of data used in determinations of 1—rates of diffusion of boron in austenite and 2—solubilities of boron in the a and phases of iron and steel&apos; has provided an equation for the diffusion of boron in a iron. In brief, the previously published data were accumulated from the results of deboron-izing (and decarburizing) experiments carried out in the range of 700° to 1300°C. Diffusion coefficients (D?) for boron in austenite were calculated using the Grube solution of Fick&apos;s law. However, only solubility values were estimated from the discontinuous concentration-penetration curves, Fig. 1, which are characteristic of diffusion in two phases. Dr. Carl Wagner&apos; has recently provided a solution for calculating D values from penetration curves of this type as follows: Cl1.1 —C8/C8- C11.1 = vp ? e ? erf (?) and D = ?/4?2 t where C11,12, C8, C11, and E have the values shown in Fig. 1. D is the diffusion coefficient, sq cm per sec; t is time of deboronizing anneal, see; and ? is a di-mensionless parameter. For a given diffusion experiment, the value of ? can be readily obtained by graphical solution from a plot of y vs the right-hand part of Eq. 1. D may then be evaluated from Eq. 2. The avplication of this solution to previously re-ported results, of which excerpts are given in Table I, permits the calculation of diffusion coefficients for boron in a iron. On the basis of these meager data, it is tentatively concluded that the diffusion of boron in a may be represented by the equation Da = 1011 e-92,900/RT where R is the gas constant, cal X 0C-1 X mol-1; and T is the absolute temperature, OK. Although the frequency factor, 106 sq cm per sec, is admittedly several orders of magnitude higher than expected, the value of Q, 62,000 cal per mol, appears reasonable and is, in fact, very similar to that for the self-diffusion of iron. It is pertinent to mention at this point that the value of Q obtained for the diffusion of boron in austenite by means of the Wagner solution is 19,000 cal per mol and is in excellent agreement with the value previously reported1 in the equation DT = 2xl0-3 e-21,000/RT [4] which was determined by the application of the Grube solution to other data. The fact that determined constants A and Q in the equation D? = A e-Q/RT were practically the same, independent of whether the Wagner or Grube solutions were used in determinations of D values, strengthened the authors&apos; belief that the computed values of D (Table I) using the Wagner solution are reliable. The relative values in Eqs. 3 and 4 for the diffusion of boron in the a and ? phases, respectively, suggest that boron forms a substitutional solid solution in a iron and an interstitial solid solution in ? iron. The same tentative conclusion has been reached by McBride et al.3 on the basis of relative solubilities, atom diameter, and the size of the interstitial hole in a and ? iron. In connection with the data for test 14 listed in Table I, it is of interest to calculate the solubility of boron in a iron using the D value given by Eq. 3. As might be anticipated from the small movement of the interface in test 14, proper substitutions in Eqs. 2 and 1 give a low value, approximately 0.0004 pct B, at 850°C. Apparently at 835°C (test 13) it is possible to obtain 0.0018 pct B, and at 850°C only 0.0004 pct B into solution in the a phase before a second phase appears. These observations are consistent with the Fe-B constitution diagram proposed by McBride, Spretnak, and Speiser.3 References P. E. Busby, M. E. Warga, and C. Wells: Diffusion and Solubility of Boron in Iron and Steel. Trans. AIME (1953) 197, p. 1463; Journal of Metals (November 1953). &apos;W. Jost: Diffusion in Solids, Liquids, Gases. (1952) pp. 69-71. New York. Academic Press Inc. "C. C. McBride, J. W. Spretnak, and R. Speiser: A Study of the Fe-Fe2B System. Trans. ASM (1954) 46, p. 499.

Jan 1, 1955

WELL known member of the Institute, Henry Krumb, survived the Japanese earthquake and has written a most interesting description of his personal experience to a friend in New York, an extended excerpt from which we publish, by permission, below. Not having been written for publication it has increased value as a human document in its vivid description of personal impressions and sensations in the midst of such a catastrophe. The earthquake was the most terrible experience I have ever passed through. We arrived at Miyanoshita, Friday, Aug. 31, about 6 p.m., having left Yokohama about 2:30 that afternoon by automobile. It was a most beautiful ride, although rather warm, but not unusually so, as it has been uncomfortably hot ever since we arrived in Japan. The first part of this automobile trip was along the old highway from Kyoto to Tokyo, which is lined with fine old pine trees. There are 52 villages along this highway and we saw real Japanese country life: naked children, both boys and girls, and men wearing nothing but gee-strings. The country on either side was intensively cultivated, the green rice fields and lotus ponds making a most beautiful sight. The occasional hills were all re-forested with young pines. In places the road runs close to the ocean beaches. The last part of the trip was up a steep mountain road, cut out of the side of the mountain. On one side of the road, in the gulch below, was a roaring river. The road crosses this river several times. Miyanoshita is beautifully located on the side of a mountain about 400 ft. above this river at an altitude of 1400 ft. The grounds of the hotel are laid out in Japanese style, small lakes with fish, waterfalls and dwarf trees making a very pretty sight. On the morning of Sept. 1, after walking around the garden and the hotel, we went to our room and I started to write up a few notes. Mrs. Krumb was reading a guide book. A few minutes before noon she came over to the table where I was writing and said that it was about lunch time. Just at that moment hell broke loose. It was the first earthquake I have ever experienced and I did not immediately realize what was happening. The building was shaking as a mine office might shake when heavy blasting was going on near by. Mrs. Krumb, having experienced many earthquakes, stepped to the open door between our room and the bathroom and called to me to get out from under the chande-lier. The plaster was commencing to fall and a large wardrobe fell over. I then knew what was happening and ran to the win dow. Before I reached it a large chunk of plaster from the ceiling hit me on the head and cut my forehead and nose. Mrs. Krumb ran for the other window just as the marble mantelpiece and the dresser fell. Without any hesitation, she jumped out of the win-dow and I followed suit out of the other. Luckily neither of us sprained an ankle-we had jumped 8 ft. We started to runaway from the building but had not gone more than 25 ft. when a second shock or rather a continuation of the first became so violent and the ground under our feet was trembling so that it was almost impossible to stand up without holding on to a tree. The ground seemed to have become liquid and rocked violently, it seemed to me in a vertical direction. In some later shocks the movement was more horizontal. About this time Mrs. Krumb saw the blood on my face and commenced to break down, but only for a moment: As soon as she saw it was not serious she was all right and during the following days stood the experiences and hardships as well as the best of them:

Jan 11, 1923