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By H. A. Kursell

IN pre-war days, the mining industry of Russia occupied in the industrial life of the Empire a place of ever increasing importance. The peak of produc-tion was reached in 1913 and 1914, the years of greatest prosperity, and a rapid growth of the mineral industry seemed assured, especially as regards base-metal mining. Exceptions were the gold production, for which a gradual decline was anticipated, and the platinum out-put which was expected to remain steady. A greater productivity of labor, corresponding to better working facilities was noticeable, and the economic status of the working classes was constantly improving. The tremendous strain on Russia's economic life caused by the World War did not materially alter the prospects of the mineral industry and the fundamental changes began only with the revolution in the spring of 1917. Their scope and character are illustrated by the figures in Table 1 which compare, for the parts of the Russian Empire under Soviet (R. S. F. S. R.) control, the average mineral production of 1913 and 1914 with that of 1920, 1921 and 1922, as supplied by bolshevik sources.1The latter three years in Russia's history were characterized by the firm establishment in power of the Communist party as a new ruling caste. Silver, lead, zinc, tin, nickel, tungsten, barite, etc., were either produced in negligible quantities, or not at all. However, before the revolution the development of several argentiferous lead-zinc deposits in western Siberia had reached a stage which seemed to assure an appreciable output of the white metals for the near future.

Jan 7, 1923

By H. R. Ogden, E. S. Bartlett, A. G. Imgram

Unmodified disilicide coatirigs were applied to sheet-tensile specimens ofCb-Dg3 and Mo-TZM veJractovy- metal alloys. Coating thickness, degree of coating-substrate interdiffusion, and specimen geonzetry (notched and plain were included in the variables studied. Tensile tests were made to determine the ductile-lo-brittle transition temperature. The disilicide coating modestly increased the transition temperatlre of TZM, but had no effect on 043. Neither material condition (recrystallized or stress-velieved) nor specimen geometry (notched or unnotched) significantly altered the effects of coatings on the transilion temperatures of. the alloys. Cracks in the brittle coatings did not propagate into the substrate, and fracture modes appeared to be the same for both un-coated and coated specimens. MOST potential structural applications for refractory metals and alloys involve exposures to oxidizing environments at elevated temperatures. The general lack of oxidation resistance of these metals will require protective coatings to allow fulfillment of their potential. Currently preferred coatings for the oxidation protection of refractory metals are brittle intermetallic aluminides or silicides. These are typically formed on the surface of the refractory-metal substrate by a diffusion reaction between the substrate and a gaseous or liquid medium that is rich in aluminum or silicon. Because of the brittleness of these coatings, they will sustain no plastic deformation at low temperatures. They are frequently cracked by cooling from the coating temperature because of the thermal-expansion mismatch with the substrate alloy. Even if they survive cooling intact, they crack rather than sustain deformation under load at low temperatures. Thus, when a coated refractory metal is strained beyond the elastic limit of the coating at low temperatures, the mechanical environment of the substrate would include both static and dynamic cracks. These might be expected to influence the flow and fracture behavior of the substrate. This could be manifested in an altered fracture mode and/or an increase in the normal ductile-to-brittle transition temperature of the refractory-metal substrate. This paper presents the results of a research program that was conducted to determine the influence of the presence of a brittle surface coating on the low-strain-rate tensile behavior of typical refractory metals at low temperatures. EXPERIMENTAL PROCEDURES Material Preparation. Thirty-mil-thick sheets of molybdenum TZM alloy (Mo-0.5Ti-O.1Zr) and colum-bium D43 alloy (Cb-IOW-1Zr-O.1C) were obtained commercially. These alloys were selected as substrate materials representing two classes of materials important in current refractory-metal technology. The TZM was in the stress-relieved condition, and exhibited a heavily fibered grain structure. The D43 had been processed by the duPont "optimum" fabrication schedule,' and exhibited slightly elongated grains typical of this process. Tensile specimens of two geometries were prepared from these materials: 1) plain specimens with 0.2-in.-wide 1.0-in.-long gage sections; 2) specimens similar to above, but with a 0.06-in.-diam hole drilled in the center of the gage section, providing a stress concentration factor, Kt, of 2.5. The "notch" geometry was selected to represent a typical condition of a rivet hole or other geometric discontinuities as might be encountered in various applications. Machined specimens were degreased, with a final rinse in acetone, prior to the application of coatings. Specimens of each substrate and configuration were pack-siliconizedin a particulate mixture of 80 pct A1203, 17 pct Si, and 3 pct NaF. Specimens were embedded in this mix (contained in graphite retorts) and coated in an electrically heated argon-atmosphere furnace under time-temperature conditions to effect nominal 1- and 3-mil-thick silicide coatings: Coating Thickness, mils Thermal Treatment 0.6 to 1.4 24 hr at 982°C 2.4 to 3.2 48 hr at 1093°C Coating kinetics were similar for both the TZM and D43 substrates. These treatments had little or no visible effect on the substrate microstructure as determined by optical metallography. The coatings on TZM were essentially single-phase unmodified disilicides, while those on D43 showed substantial evidence of modification by proportionate reaction with the respective substrate elements or phases, as shown in Fig. 1. It was recognized that these coatings might not be particularly desirable regarding protective capability. However, it was desired to circumvent possible inter -ferring chemical interaction with the substrate by pack additives such as chromium, titanium, boron, aluminum, and other elements that typify the better protective coatings for these materials.' Thus, the results presented apply specifically to the simple silicide coatings investigated. They may not be rep-

Jan 1, 1967

By A. L. Holley

THE occurrence of this gas, in quite appreciable. quantities, has been observed for many years in its escape along the creeks of Western Pennsylvania, and more recently, in much larger quantities, from the oil wells in and near Venango County. Within about five years several gas wells have been struck, from which the gas has escaped continuously in prodigious quantities and at great pressure. During the last year a systematic attempt has been made to utilize the flow of gas from one of these wells, and results have been reached which were quite unexpected. This well, near Leechburg, Armstrong County, is about twenty miles northeast of Pittsburgh, on one of the side tributaries of the Alleghany River. It had been drilled in search of oil to a depth of 1250 feet in 1871, but none was found. A great flow of gas was developed, however, accompanied by a slight spray of salt water, and this has continued with little or no diminution to the present time. The gas in its escape has been discharged through a 5-inch pipe, and at a pressure of from 60 to 80 lbs. per square inch, but no exact determinations, either of the pressure or quantity, have been made. The rolling mill of Messrs. Rogers & Burchfield is on the opposite side of the river, and it has been for some years devoted to the production of fine grades of sheet iron from charcoal pig metal, by puddling and in knobbling fires. The usual weekly product of the mill has been 30 tons of No. 30 tin plates, and 50 tons of No. 24 to 28 sheets. The well was bought by this firm for $1000, and the gas is led across the river, a distance of about 500 feet, through a 3-inch pipe. It is distributed through half-inch pipes, and at a pressure of about 45 lbs. per square inch, to several of the furnaces, but a large fraction of the total flow of gas is still wasted at the well. No essential alteration in any of the furnaces has been found necessary in the use of the gas fuel, except to brick up the fire-bridge and to put in the gas and air pipes. The old grate used for coal is loosely covered with bricks and cinder, so that a slight percolation of air may take place through them. The gas is admitted through a half-inch pipe, and blows toward the fire-bridge through 18 or 20 one-eighth inch jets. The air is blown in at about 2 lbs. a

Jan 1, 1876

By Ralph H. Sweetser

In these days of the almost exclusive use of byproduct coke as the blast-furnace fuel in this country, it may seem out of place, and smacking too much of reminiscing, to say anything about the use of anthracite coal for blast-furnace fuel. Viewed in the light of present day research and of developments in alloy steel, some observations seem appropriate regarding this raw material that played such an important part in the development of the iron and steel industry during 50 years of the past century. The research work of that period, in which the whole iron and steel industry was revolutionized, passing from wrought iron to steel and from charcoal to three new mineral fuels for blast furnaces, was done chiefly in the full-sized plants of that day; and it was done largely by the ironmasters themselves. The laboratory had not become a part of the ironworks, and scientific instruments were lacking around the blast furnaces; "rule o' thumb" and erroneous theories prevailed—and have not yet all been eliminated. The recent interest in meteoric iron in connection with nickel steels and other alloys proves that it is worth while to search for the reasons for excellence and the causes for failures in the metallurgy of bygone days. Turner1 gives the reasons why the blast furnace of Dr. F. W. Geissen-hainer on Silver Creek, Schuylkill County, Pennsylvania, about 10 miles northeast of Pottsville, was operated successfully with anthracite2 in August and September, 1836, and why several other attempts failed. It has taken a long time for us to find out that there are great differences in the behavior of anthracites from the several parts of the anthracite region, as well as differences in coking coals from the bituminous fields.

Jan 1, 1935

ONE of the largest hard rock operations ever undertaken on the continent is part of the Aluminum Co. of Canada's gigantic British Columbia development in the mountainous wilderness back of the north coastal port of Prince Rupert. To provide electric power for the huge aluminum smelter under construction at Kitimat, 400 miles northwest of Vancouver, Alcan is reversing the flow of a 5000 sq mile drainage area by blocking the eastward-flowing Nechako River with the largest sloping clay core dam in the world. The backed-up waters of a 110-mile chain of lakes will be tapped at their eastern end by a 10-mile tunnel piercing a mile-high mountain range. From this tunnel water will pass through penstocks to a subterranean powerhouse after a drop of 2600 ft and empty into the Kemano River almost at tide water. A projected battery of eight turbine-driven generators will produce more than a million horsepower, although, initially, three generators of 140,000 hp each will be installed. Power will travel 50 miles across another mountain range, over a glacier, and across an arm of the ocean to what will ultimately be the world's largest aluminum smelter-a plant 2 ½ miles long. Ore smelted at Kitimat will come from Jamaica, but carving the tunnel, penstocks, and powerhouses cavern out of the heart of a mountain is so vast a job that most of the problems of a major mining operation are encountered. The powerhouse cavern alone is a big hard rock job. This chamber lies 1400 ft back from the mountain face and will eventually be 1100 ft long, 80 ft wide, and an average 118 ft high. Now under construction, the roof of this cavern will take the form of a parabolic arch. The prime contractors, Morrison-Knudsen Co. of Canada, are to excavate a length of 700 ft by May 1954. This section will provide a main control room, erection bay and space for eight generating units. The roof arch is to be lined with reinforced concrete and a plan now under study provides for concrete curtain walls separated from the rock sides of the cavern to prevent moisture accumulation. Cavern Excavation Excavation of the cavern began with the driving of an 8x8-ft exploratory drift 1500 ft long. Two horseshoe-shaped access and tailrace tunnels were bored to serve excavation work. These tunnels are approximately the same length as the exploratory drift and 25 ft high. The floors of both tunnels were graveled to permit the use of rubber-tired equipment. Eventually, the tailrace tunnel will be enlarged to 40 ft height to drain water from the turbine pits below the powerhouse floor. When the tunnels reached the cavern location, a mucking tunnel, approximately 25 ft square, was driven the projected length of the powerhouse at the floor level of the cavern. Another tunnel, 10 ft square, was carried in the same direction from the exploratory drift at the level of the arch spring line, 65 ft above the cavern floor. To connect these two tunnels, seven shafts 10 ft wide and approximately 80 ft long were driven through to the lower mucking tunnel, 65 ft below. These shafts were located on 120-ft centers, the length of the powerhouse cavern. Blasting is done from this upper tunnel to the final line of the roof arch. Two bulldozers, an International TD-24 and a TD-9, doze blasted material to the shafts where it drops to floor level. From here, the rock is loaded by Eimco loaders and power shovels into rubber-tired Koehring Dumptors and Euclid end-dump trucks. The arched roof section will be completely excavated and concreted before vertical blast holes are drilled down from the level of the arch spring line to the floor of the cavern. The arched roof section with its parabolic curve will have a width of 80 ft and a rise to the crown of 47 ft.

Jan 1, 1952

By Roy Lindgren

Since the year 1643, when the first blast furnace in America for treating iron ore was built at Saugus, Mass., out of mica schist quarried in the neighboring district, the procurement of a suitable refractory for furnace lining has been a problem of concern to the operators of furnaces. The stacks built of mica schist continued to smelt iron ore until about 1836, when, according to F. H. Norton, the first firebrick were produced1, at Queens Run, Pa. Other writers speak of brick having been molded and burned in Massachusetts about the year 1834. In 1841, Andrew Russell began to produce medium refractory plastic clay brick near East Liverpool, Ohio, that were used for lining blast furnaces1. The well-known Kentucky clay-producing district was not opened up until the year 1871, but since then it has produced a large percentage of the linings for iron blast furnaces. While some strides have been made by the refractories industry during the 100 years that have passed since the first firebrick were produced, it has been only during the last two decades that any real progress has been made towards bettering the product, even though the method of production had improved. Perhaps the fault lies with the user of the brick rather than with the producer, for not sooner demanding a supe¬rior product. During the past 15 years the tonnage produced per lining has increased from 500,000 gross tons to 1,000,000 gross tons, and now some furnaces are producing 1,600,000 gross tons and better on a single lining. It is true that enlarged capacity of furnaces and improved practice have accounted for some of this increase in tonnage, nevertheless better quality in firebrick must be given credit for its share. However, we are not yet ready to say that we have reached a maximum life of furnace lining. We believe that a better product can be produced and that the refractories industries of America will, through their extensive research depart-

Jan 1, 1936

By N. J. Grant, D. S. Bloom

An investigation of the Cr-Mo-Ni ternary system and attendant binaries has been completed. Some changes in the binary diagrams are shown to be necessary. The 1250°C section of the ternary and the liquidus surface have been delineated. A system of invariant reactions existing above 1250°C has been suggested. BECAUSE chromium, molybdenum, and nickel are important not only as bases for alloying systems but also as major alloying additions for other refractory metals, a study of the phase diagrams is of great interest and importance. In addition there are interesting features of theoretical value, a number of which are of prime consideration in this work. The major aim was to determine the ternary phase diagram of these metals at and above 1250°C, with sufficient attention given to lower temperatures to permit an understanding of the conditions existing at 1250°C and above. To understand ternary systems it is essential that accurate information concerning the binary systems be available, which in turn presupposes thorough knowledge of the constituent elements. On the basis of this investigation not only were significant inaccuracies indicated in the binary systems but much consideration had to be given to one of the elements, that is, chromium. The observation has been made that chromium undergoes a crystal transformation at a temperature near its melting point. Much of the experimental data leading to the detection of this transformation and the effects of this hypothesis on the form of the Cr-Ni phase diagram have already been reported' . and will not be repeated here. The experimental methods employed in the investigation have also been covered.'-" he purity of the molybdenum utilized in this work showed by qualitative spectrographic analysis that chromium, copper, iron, silicon, and tungsten were present in quantities of 0.1 to 0.001 wt pct. During the course of the investigation more than 100 melts were made, weighing from 150 to 200 grams each. Each of these melts was sectioned for chemical analysis and for specimens for examination in the as-cast and various heat-treated conditions by X-ray and metallographic means. All X-ray diffraction patterns reproduced here were made using chromium Ka radiation although iron and copper targets were used on occasions during the investigation. Experimental Results Mo-Ni-System: The most recent comprehensive investigation of the Mo-Ni system was made by Ellinger' in 1942; the diagram as shown in the American Society for Metals Handbook is based on his work. The work done on the system in this investigation substantiated the general form of the diagram although some differences were observed. These are indicated by a comparison of Ellinger's diagram and Fig. 1, the latter showing the results of pertinent thermal analyses made in this research. It can be seen in Fig. 1 that the eutectic composition is

Jan 1, 1955

By W. M. Lehrer

RECENT development of high-temperature brazing alloys has required information regarding the eutec-tic temperature and composition of the Pd-rich section of the Pd-B system.. Literature on this subject was found to be scanty with no complete system yet determined. Buddery and welch1 report an inter-metallic compound occurring in the Pd-B system at 6.33 wt pct B (PdsBz). Another source, Sieverts and Bruning,2 indicates that after annealing at 700°C, a 1.6 and a 2.0 wt pct B alloy were two phase whereas a 0.75 pct B alloy was single phase. Recent work3 using a hot-stage microscope has proposed a eutectic temperature of 743OC. This method utilized the observation of the onset of incipient melting at grain boundaries. The eutectic composition, however, was not determined. METHODS AND PR0CEDURES Five alloys were prepared by triple inert arc-melting high-purity palladium (95.98 pct) with additions of 2 to 6 pct purified amorphous boron. The alloys were crushed within a hardened steel mortar with a steel pestle. The 2 pct alloy, however, was too tough to be crushed in the above manner so was crushed by successive reductions through a rolling mill. Approximately 17 to 18 g of the crushed alloy was used in the thermal analysis. The alloys were melted in pure quartz crucibles. The thermocouple protection tubes were also of quartz and fused along the inside wall of crucibles to prevent floating of the thermocouple. This was an expediency as centering of the thermocouple in the melt would have been preferable. Melting was done in a tube furnace with a slow flow of tank helium maintained as a protective atmosphere. Heating and cooling rates were maintained at 3.5OC per min through the critical temperature ranges by means of a synchronous motor hook-up to a furnace controller. All specimens were heated to 100°C over the indicated liquidus on heating prior to reversing to the cooling cycle. Specimens emerged from the furnace with mirrorlike brightness attesting to the protective quality of the helium atmosphere. Thermal arrests were detected by the differential thermocouple technique using chrome l-alumel couples and nickel as the reference specimen. Arrests were recorded on a Leeds-Northrup automatic X-Y recorder. The specimen temperature and temperature difference were read to + 7.5OC and ± 1.5OC, respectively. These readings were well within the sensitivity of the instrument amplifier. Chromel-alumel couples were used in lieu of Pt-Rh as very few arrests were contemplated in excess of 1200" to 1300°C, the approach of instability. No standardization was performed on these thermocouples subsequent to their use. RESULTS The thermal arrest temperatures obtained are plotted in Fig. 1 as a function of wt pct B. The composition function was obtained by chemical analysis of portions of the specimens after the thermal analysis was completed. The 2 to 6 pct nominal alloys were found to have 1.80, 2.35, 3.65, 4.18 and 5.90

Jan 1, 1960

By J. B. Mackintosh

The general method which has been followed since the time of Heinrich Rose, and perhaps before, for the determination of phosphorus in iron and steel, is to dissolve the sample either in nitric acid, aqua regia, or some other oxidizing agent, in order to convert the phosphorus into phosphoric acid. In 1832 Rose* recommends HNO3, or aqua regia, preferably the former. Wohler, in 1854,† directs that the iron should be oxidized by burning in a stream of oxygen for the determination of carbon, and the resulting oxide of iron containing all the phosphorus as phosphoric acid be treated as an ore. In the fourth English edition of Fresenius' Quantitative Analysis,† pp. 666-667, we find directions to dissolve the iron in aqua regia or bromine; and, coming down to more recent times, we find F. A. Cairns § directing the employment of a boiling solution of potassium chlorate and hydrochloric acid for the oxidation and solution of the sample to be analyzed. It is unnecessary to multiply quotations from the literature of the subject, though they might be extended almost inclefinitely, since it would only be a repetition of the above mentioned series of methods in one form or another; and the references given have been selected to show how little essential variation there is in the type-methods of solution used. In fact it seems to be universally accepted as a necessary first step in the determination that some such method of solution should be followed. One chemist, however, struck out on a different line of thought, and to the Lest of my knowledge only one, M. V. Tantin, who, in

Jan 1, 1886

By I. S. Servi, N. W. Marr

TIN markedly increases the softening temperature of pure copper with only a moderate effect on conductivity. Smart and smith' indicated a substantial increase in softening temperature upon addition of different levels of tin up to 0.05 pct in oxygen-free copper. However, after finding no effect of tin in copper containing about 0.025 pct O, they concluded that "the presence of oxygen completely nullifies these changes in properties by precipitation of SnO2.'' Since preliminary tests on Cu-Sn-0 alloys indicated that this conclusion is not valid outside the extremely limited tin range investigated by Smart and Smith, research was primarily directed toward the partition of tin between the copper-base matrix and the oxide phases. Experimental. Material was prepared by sealing an 80 to 100-g charge consisting of oxygen-free copper, tin, and copper oxide into a y-in.-ID quartz capsule which was evacuated and back-filled with tank-grade argon to 3 atm. The sealed vials were heated to 1125ºC, kept at temperature for 30 min, occasionally agitated, and then allowed to air-cool outside the furnace ("as-cast"). The castings were cold-swaged to 0.370 in., annealed 1 hr at 600°C ("as-annealed"), and cold-swaged 16.5 pct ("as-swaged"). Cross sectional wafers were annealed for 1 hr at various temperatures and Rockwell F hardness measured. The "softening temperature" was defined as the interpolated temperature at which the hardness was 50 Rockwell F. Electrical resistance was measured by an ac poten-tiometric method on 0.040-in.-diam wires after a 1 hr anneal at 600°C. Relative accuracy was of the order of 1 pct. Chemical analyses for tin were performed by two methods, identified as "Acid flux" and "HN03 dissolution". The first consisted of fusing the sample with potassium pyrosulfate and leaching the resulting melt with 10 pct sulfuric acid. The spectrophotometric procedure which followed is described in Ref. 2. Oxygen analyses were performed by the vacuum-fusion method with a bath temperature of 1650°C. Tin content of oxides and copper phases was measured with the X-ray microprobe, using a calibration curve based on a series of oxygen-free Cu-Sn alloys. Results. In Fig. 1, the nominal tin composition for alloys containing nominally 400 ppm O was plotted vs the analyzed or the computed tin content—the latter derived from one of the following: 1) "acid flux" chemical analysis; 2) "HN03 dissolution" chemical analysis; 3) X-ray microprobe analysis of the copper matrix; 4) equivalent tin content, based on the measured increase in electrical resistivity (an increase in resistivity of 1.6 microhm-cm is approximately equivalent to 1 wt pct of sn1,3). It is evident from Fig. 1 that the results of the acid flux analysis are consistent with the nominal composition while the HNO3 dissolution method generally yields lower values, consistent with those obtained from electrical resistivity measurements and somewhat higher than those obtained by microprobe analysis. Microprobe and resistivity values are affected only by the tin in solid solution; hence the HN03 method

Jan 1, 1970

By V. Runolinna, R. T. Hukki

Preliminary separation of precious metals by flotation can offer a simplification of the conventional method of treatment of anode slimes. Laboratory flotation experiments show that rich gold and silver concentrate can be obtained with excellent recoveries. THE purpose of this paper is to present the results of an investigation into the possibilities of separating precious metals from anode slimes by selective flotation. The work was carried out at the State Research Institute in Helsinki, Finland. The sample of anode slimes tested was submitted by the Outo-kumpu Co. in Finland. In 1947 the average assay of filtered and dried anode slimes in the Outokumpu refinery in Pori was as follows: Au, 0.50 pct; Ag, 9.38 pct; Cu, 11.02 pct; Ni, 45.21 pct; Pb, 2.62 pct; Fe, 0.60 pct; Sn, 1.00 pct; Sb, 0.04 pct; As, 0.70 pct; Se, 4.33 pct; S, 2.32 pct; SO,, 2.17 pct; SiO2, 2.25 pct. The daily assay naturally varies considerably from the figures presented above. The present conventional method of treatment of anode slimes consists of the steps shown in the flowsheet in fig. 1. As seen from this simplified flowsheet, the present method of treatment involves many steps by which the components of the anode slimes are removed one after another until the main products, gold and silver, may be separated. Due to the small scale of operations, the above procedure is carried out in small batches requiring considerable amount of manual labor. If, however, the anode slimes could be treated in

Jan 1, 1951

By Bradley Stoughton

FRom 1886 to 1889, approximately, Henry M. Howe was engaged most actively in preparing his Afetallugy of Sieei, which became the foundation for his leadership in the science of physical metallurgy of steel. During the same period, Hadfield was developing the first of the industrially-important alloy steels, and Riley was initiating the commercialization of nickel steel. Both of these alloys are treated briefly in Howe's book; indeed, he missed practically nothing of interest in all the steel literature preceding 1889. There had been many good books on steel before his, but none of the same character. His genius, like that of Aristotle and Francis Bacon, manifested itself in the painstaking collection of data, followed by their coordination and interpretation. Not only manganese and nickel steels were discussed, but several which most of us, with our tendency to overlook the past, are wont to regard as exclusively recent developments, such as: silicon, chrome, copper, tungsten, vanadium, titanium, chrome-tungsten and many other steels. But Howe never devoted much special attention, in his writings, to alloy steels; the science of the relations of iron and carbon offered field enough even for his capacity for constructive thought. And, today we find the science of alloy steels in much the same condition as was the metallurgy of steel before Howe wrote his first great book. Many, many important facts have been presented to the world, especially by Sir Robert Hadfield, James Riley, Henri le Chatelier, Kotaro Honda and his associates, Maunsel White, Harry Brearley, and a host of others. Several books have been devoted wholly or chiefly to alloy steels, such as those by Hibbard, Mars, Heyn and Wetzell, Grube, Burnham, Gillett and Mack, Billón, Brearley, Monypenny, Becker, Savoia, Sauveur, Guillet, Tammann, Guertler, Goerens, Guia, Wuest, Ledebur, Schenck, Aitchison, Osann, Guettier, Buchetti, Escard. The vast number of data on the subject are indicated by a bibliography' which was collected in preparing this address. It covers the years 1911 to

Jan 1, 1927

By A. B. Kinzel, Walter Crafts

Inclusions of nonmetallic matter have long been recognized as objectionable in steel. A complete theory of the effect of inclusions, which is consistent with that held today, was outlined in Howe's early work on the metallurgy of iron and steel. This, by analogy with wrought iron, assumes that the observed difference in properties between specimens with and across the direction of rolling is due to slag enclosures. The fact that wrought iron shatters in ballistic test is taken as an indication of the effect of inclusions on the resistance to'shock. Since that time, the development of the metallurgical microscope has shown the discrete nonmetallic inclusions in the iron matrix. Better resolution has brought out the great number of inclusions, and the availability of good microscopes in steel mills and laboratories has constantly drawn more attention to these inclusions. This recent emphasis on inclusions has resulted in studies as to their nature, mode of formation and methods of elimination, such as the work of Wohrman, Dickenson, Herty and Schenk. In addition to the published works, many investigations which have not been published have been carried out by private laboratories and have resulted in the application of specifications intended to limit inclusions in many qualities of steel. Neverthless, a clear understanding of the quantitative effect of inclusions on the physical properties of steel has been lacking. It is the specific purpose of this study to determine the quantitative effect of inclusions on the physical properties of steel that determine its behavior under dynamic stress. The study has shown that a quantitative relation between counted inclusions and impact strength does exist in normalized steels, but that factors other than readily visible inclusions are equally important. Furthermore, the present commercial rating of inclusions is unreliable with respect to counted inclusions or dynamic tests, and this study indicates that the tensile-impact test is more reliable as a measure of quality. General Scheme The general plan of the study follows: (1) To develop or establish a satisfactory measure of dynamic strength; (2) to develop or establish a

Jan 1, 1931

By Erhard Hornbogen

Twins were propagated into large, well-annealed crystals of a, iron-phosphorous and a, iron-molybdenum solid solutions. Strain fields caused by interaction of these twins were made visible by precipitation. These strain fields can be explained by assuming that transmitted and reflected plastic waves are created when a twin strikes an obstacle with high velocity. The twin front itself can be regarded as a shear wave that may be able to develop a one-dimensional shock front. The occurrence of {WO} fracture by the reflected stress is also discussed. The described effects were not found when twins propagated at lout velocity into disturbed crystals, for example, into crystals previously deformed by slip. DEFORMATION twinning leads to shearing of a crystal lattice under external load. The shear appears to be homogeneous over a large number of crystal layers. In a iron, shear occurs on the (112) planes and in the <1ll> directions. The formation of a twin proceeds by the movement of partial dislocations that change the stacking at the twinning plane.l, In the bcc lattice, this partial dislocation can be created by dissociation of a sessile a/2[111] dislocation that lies in the (112) plane: Cottrell and Bilby&apos;s spiral mechanism can explain the homogeneous shear, but has not been proven experimentally. Nevertheless, this mechanism can explain the possible high velocity of formation of a twin. In Cottrell and Bilby&apos;s model, as in other models of twinning or phase transformation, a much higher energy is required to spread the first stacking fault than to add new layers by further movement of the partial dislocations. The dislocations will therefore become accelerated proportional to the difference between the stress, tN, to nucleate and tp to propagate the twin. It has been assumed3,4 that tn is the stress required to make two partial dislocations of apposite sign pass for the first time. After the first layer of new stacking has been formed, the maximum stress is TN=Gb/4pd where G is the shear modulus, b is the Burgers vector of the partial dislocation, and d is their perpendicular spacing so that the amount of shear is S = b/d. From this formula, a value of 25 kg per mm2 has been estimated for twinning in Zn3. The measured values are smaller, but local internal stresses may help to provide the necessary stress. For a iron, a higher value can be expected, mainly due to the higher value of G. That agrees with the observation that "burst" twinning occurs if the yield strength is in-

Jan 1, 1962

By E. A. Sondreal, F. I. Honea

This paper summarizes coal and ash properties and pilot-plant results for 14 key lignite seams in the Gulf Province, as well as discusses operating problems encountered by the four Texas power plants burning Gulf Province lignite. Some consideration is given to the design of future power plants.

Jan 1, 1982

By Frank E. Hall

The needs of the World War showed the necessity of a metal stronger than cast iron which would supplement the supply of steel. SO patriotic metallurgists were spurred to new efforts to improve the status of that half-breed of the metal world that had been, more or less erroneously, christened semisteel. As a result, it was developed that for certain purposes semisteel was superior to steel in the manufacture of shells. Steel, because of its greater strength and toughness, is able to carry a greater charge of explosive, weight for weight; consequently, a steel shell can be expected to cause greater destruction when used against fortifications or other inanimate objects, while semisteel, because of its comparative brittleness and its resulting greater fragmentation, is capable of inflicting more serious loss of life when used against opposing forces of men. Therefore, each has its proper place in the economy of war. For the present purpose, steel may be defined as a comparatively pure iron whose physical nature is modified by small amounts of carbon, silicon, sulfur, phosphorus, manganese, and possibly other rarer elements; while cast iron generally is limited to approximately 93 per cent. of iron, the remainder consisting of larger percentages of carbon, silicon, sulfur, phosphorus, and manganese. Semisteel may be defined as cast iron in which the carbon and silicon have been reduced to as low a point as is consistent with maintaining the nature of gray iron, and in which the remaining elements are adjusted so as to give the greatest strength. The fact that the addition of steel scrap, in varying percentages, is the means used to reduce the carbon, together with the fact that a considerable increase in strength results, is probably the explanation of the origin of the name semisteel. That ordinary grades of steel show from three to four times the tensile strength of cast iron is due to the fact that the crystals, or grains, of cast iron are interspersed with flakes of

Jan 1, 1922

By Charles F. Park

Mining in the Pacific Coastal Region has passed through three stages of development. First came the gold rush days, a period when gold and silver were the objects of intensive search. Second was the development of the high-grade base-metal deposits and of the smaller and more refractory precious metal deposits. And third was the modern period that covers the discovery and development of the rarer metals, of large low-grade deposits, and of the industrial and nonmetallic minerals. Brief discussions of the salient geologic features of the region and of a few of the better known mineral deposits are given. The regional geology is favorable for continued exploration and discovery of new deposits and districts.

Jan 1, 1968

By Max Roesler

I. INTRODUCTION THE following article concerns the geological occurrence of the iron-ore deposits on the south coast of Cuba. The article is based on a detailed field study, made in the hope that some information would I gained which might be of value in the search for further orebodies, Or in the economic development of the ore already found. LOCATION The iron-ore deposits of the Firmeza district lie near the coast, in the southeastern part of Cuba, in the Province of Oriente. The orebodies of this district form part of a belt of deposits that extends from Sigua, 25 miles east of Santiago, to Sevilla, 5 miles east of Santiago, and lies on the seaward slope of the Sierra Maestra range of mountains.

Jan 10, 1916

By S. R. Seagle, R. Schuhmann, N. A. Parlee

Rates of CO evolution and CO absorption were measured for liquid-iron alloys containing from 0.15 to 4.4 pet C, using a modified Sieverts apparatus. The alloys were held in alumina crucibles, so that both crucible-metal and gas-metal reactions occurred simultaneously. The data are interpreted on the hypothesis that the C-O reaction rate is controlled by oxygen diffusion across the boundary layers at the gas-metal and crucible-metal surfaces. CARBON-monoxide evolution from liquid Fe-C-0 alloys is a key reaction in steelmaking processes, both in the steelmaking furnace and in the ingot mold. Also, this reaction probably is an essential step in the desulfurization of iron under blast-furnace conditions, for which carbon is the principal reducing agent or deoxidizer present in the iron. The thermodynamic properties of the solutions of C and 0 in Fe and the equilibria of these solutions with gaseous CO and CO, have been investigated in some detail and are reasonably well understood. However, while these equilibrium data have accumulated in the laboratory, other data have accumulated to show that equilibrium conditions are often not achieved under plant operating conditions. Thus, an understanding of the rate and mechanism of the carbon-oxygen reaction not only has theoretical interest, but ultimately may assume considerable practical importance. Reaction Mechanisms and Rate Equations The over-all reaction under consideration is C (in liquid Fe) + 0 (in liquid Fe) + CO (gas) [I] Although this appears to be a straightforward heterogeneous reaction, involving just two phases (liquid metal and gas), a wide variety of reaction mechanisms has been proposed. Different authors have expressed divergent views as to the nature of the rate-controlling steps. Accordingly, a brief discussion of the possible reaction steps will be given. 1) Homogeneous Reaction Within Liquid Iron— Feild&apos; and Jette2 have applied chemical reaction-rate theory as if the C-0 reaction were a homogeneous, second-order reaction occurring within the liquid-metal phase. This mechanism necessarily produces CO "molecules" dissolved in liquid iron. In order to yield CO gas, the dissolved CO must nucleate CO gas bubbles, and the dissolved CO also must be transported to the bubble surface where it enters the gas phase. However, this mechanism is inconsistent with present concepts of the structure of liquid-metal solutions. 2) Reaction at Gas-Metal Interface—At the surface of contact between gas and metal, a second-order reaction between dissolved C and 0 may occur to produce a CO molecule which enters the gas phase. This reaction presumably involves an "activated complex" structure in the interface. Some such surface mechanism appears essential because the carbon and oxygen do not occur in the same species in the two phases and therefore must react in some way at the surface. At liquid-iron temperatures, calculations based on reaction-rate theory indicate that the surface reaction must proceed extremely rapidly and thus cannot be a rate-determining step in the C-O reaction under ordinary conditions.3 ccordingly when the over-all reaction is measurably slow, the surface reaction may be considered at equilibrium; that is Cc* Co* = m&apos;Pco [2] in which Co* and Co* are the concentrations of carbon and oxygen (weight per unit volume of metal), respectively, in the liquid-metal phase at the interface, Poo* is partial pressure of CO in the gas phase at the surface (atmospheres), and m&apos; is the mass-action constant for Reaction [I]. Since activities of C and 0 are not strictly proportional to concentrations, m&apos; is not a true thermodynamic equilibrium constant but varies with carbon content.&apos; 3) Mass Transport—If the reaction is to proceed at the gas-metal surface, as described above, dissolved carbon and oxygen in the liquid iron must come to the surface and gaseous carbon monoxide must move away from the surface. Details of the mass-transport mechanisms depend on the kind of system under consideration. For a stirred metal bath in contact with a CO gas phase, the transport of C and 0 to the surface can be described by 1 Do Mols CO evolved per sec = — 1/16— Do/do Agm (CO-Co*) = 1/12 Do/do Agm *(Co-Co*) [3] In this equation, Do and D, are the diffusion constants for oxygen and carbon in liquid iron, 6, and 6,. are the boundary-layer thicknesses, A,,,, is the

Jan 1, 1959

By R. M. Rose, D. P. Ferriss, J. Wulff

Single crystals of tantalum were grown by multipass zone melting in an electron beam apparatus. Tension tests of these crystals at 4.2 x 10&apos;* per sec-gave a resolved louler yield stress of 7300 psi at with a chisel edge; total uniform deformation increased from 1 to 2 pct at 77°K to over 20 pct at 300°K. This, and slip traces, indicated that deformation is more inhomogeneous at lower temperatures. Deformation occurred exclusively by slip in <III> directions, on planes in the <111> zones. Cross slip was frequent, particularly for those orientations where the resolved shear stress on the cross system was comparable with that on the primary plane. The existence and magnitude of the upper yield point were found to be sensitive to orientation. This effect was related to the circumvention of dislocations which were "locked" by interstitial atoms, by double cross slip, or composite slip, of screw dislocations. Prestrainirg at 300°K did not remove the upper yield point in subsequent tension at 77°K. Also, this treatment was found to cause yield points, where none existed for the unprestrained material. This behavior was also rationalized in terms of the circumvention mechanism. The mechanical properties of polycrystalline tantalum have been investigated by several workers,1"4 and may be summarized by pointing out the following features: a) Sharp rise of yield stress with decreasing temperature below room temperature. b) Response to strain aging at 300" to 400"C. c) Yielding with upper and lower yield-point phenomenon at and below room temperature. d) Low strain hardening below room temperature. e) Reluctance to twin except at very high strain rate. Furthermore, tantalum is ductile at ordinary strain rates as low as 4.2"K.&apos; This feature, in particular, makes tantalum unique among the high melting point, bee transition metals of Groups V and VI in the periodic system. Therefore, a more fundamental and detailed knowledge of the low-tempera- ture deformation of tantalum is required. Since high purity and the absence of grain boundary effects make such a study most productive and unambiguous, the method of electron beam zone melting was selected for the preparation of test specimens. This technique has the further advantage that crystals may be seeded in order to obtain any desired orientation. EXPERIMENTAL The electron beam zone-melting equipment was constructed as a modification of that described by Calverley, Davis, and Lever? All crystals were made by three passes of the molten zone at 4 mm per min to achieve an optimum purity commensurate with economy of time and material. Table I reports the chemical analyses of the starting and refined metal. It will be noted that there remain some 50 to 70 wt ppm of interstitial elements, while all substi-tutional elements have been removed by evaporation. Tungsten is picked up by evaporation of the cathode filament wire. Crystal seeding was normally successful within 3 deg, and the orientation of all crystals are given in the standard stereographic triangle in Fig. 1. Tensile bars were prepared from the crystals by centerless grinding to the shape in Fig. 2. Specimens were then electropolished to remove 10 pct of the diameter in order to remove worked material and to prepare a metallographic surface for subsequent deformation trace observations. All specimens were held in the split bushing grips as indicated in the cross section view of Fig. 2. All tests were made at a strain rate of 2 112 pct per min in an Instron testing machine.

Jan 1, 1962