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By C. Altstetter, Y. Huang, E. de Lamotte

Nitrogen was introduced into pure niobium (colutn-bium) and dilute Nb- Zr alloy wires by equilibration with pure nitrogen gas at high temperatures. Room-temperature hardness was correlated with the alloy composition, heat treatment, and microstructure. Hardness was found to be most sensitive to the amount of nitrogen in solution; however, a significant hardening up to 500 Dph look place upon aging of quenched alloys. A coarse dispersion of nitride did not significantly increase the hardness of the material though the nitride was found to have a hardness of 1400 Khn. BY now the marked effect of nitrogen on the properties of bcc refractorv metals is we ll-documented. There are few studies, however, in which the state of the solute has been carefully controlled. Indeed, the composition-temperature-pressure relationships in nitrogen-refractory metal systems have been determined precisely for only niobium and tantalum.'-4 Most investigations on niobium have dealt with alloys High-nitrogen alloys cleaved on (100) planes and had negligible ductility. Precipitation was found to occur in either a discontinuous or a continuous fashion, depending on the composition and cooling rates. Upon cooling, the precipitation first proceeded discontinu-ously by the nucleation of nilride needles at subbound-aries and subsequent growth of a two-phase microcon-stituent. At lower temperatures continuous precipitation occurred in the subgrain interior ahead of the discontinuous precipitation interface. containing nitrogen, carbon, and oxygen in significant but uncontrolled and, in several instances, undetermined amounts. Thus, there has been some uncertainty regarding the presence and the distribution of precipitates. This paper presents results which relate the hardness and microstructure to the heat treatment and nitrogen content of niobium and Nb-Zr alloys. Niobium rod* produced by the Wah Chang Corp.

Jan 1, 1968

By Albert Sauveur

The following propositions and corollaries are intended to present, as concisely as possible, some of the evidences gathered while studying the microstructure of steel. Each proposition is accompanied by photo-micrographs showing the structure of illustrative samples. But it must be understood that the special instances thus chosen for illustration are by no means the only ones on the examination of which the propositions are based. On the contrary, they have been taken almost at random among a great many, in some instances more than a hundred, similar cases. Proposition I. A slow and undisturbed cooling from a temperature x or higher produces crystallization. Plate L, Plate II., Fig. 1, and Plate V., Figs. 1, 2, 3 and 4 show the crystalline structure of several pieces of steel. Plate I. reveals this structure in the center of a slowly-cooled Bessemer ingot, 1.5 feet from the top. Plate II., Pig. 1, is from a small 2-in, square testingot of the same heat, slowly cooled. Plate V., Figs. 1, 2, 3 and 4 show the structure of different parts of steel rails. The original photographs were all magnified 70 diameters, and have been reduced in engraving to 63, or to 55 diameters, as shown in the plates. Proposition 11. Undisturbed cooling from an initial temperature lower than x is not accompanied by crystallization. Plate III., Fig. 5, shows the structure of a piece of rail-steel forged from a white heat until it had fallen to dull-red. The structure is almost amorphous. Practically no crystallization has taken place during the subsequent cooling. Corollary.—Pieces of steel finished at a temperature lower than x do not take a crystalline structure.

Jan 1, 1894

By T. D. Yensen

I. INTRODUCTION IT is safe to say that of all the different materials that go to make up electrical machinery, iron is the most important. Upon its -magnetic and electrical quality depends not only the efficiency of the machine but, to a very large extent, the cost. Aside from the mechanical strength, there are three characteristics of iron that are of importance in this connection, viz., magnetic hysteresis, electrical resistance, and magnetic permeability. In the homopolar machine, where the magnetic flux is unidirectional, permeability is the only important factor, but in all other machines the flux is reversed in one or more parts and causes losses due to hysteresis and eddy currents. The eddy-current loss depends largely upon the electrical resistance of the iron; the higher the resistance the lower the loss, under otherwise identical conditions. These hysteresis and eddy-current losses cause heating of the machine, and as they depend upon the magnetic density in the iron, this density must be regulated so that the heating shall not be excessive. If the hysteresis and eddy-current Josses and the permeability of the magnetic circuit will permit an increase in the magnetic density, or, in other words, a decrease in the amount of iron, this in turn will mean a decrease in the amount of copper needed, with a consequent decrease in the copper loss. It is thus seen that; indirectly as well as directly, the above characteristics have an important bearing on both the efficiency and the cost of electrical machinery. In general, it may be said that iron for electrical machinery should have the following characteristics: (1) low hysteresis loss, (2) high electrical resistance, and (3) high permeability. These are the characteristics that investigators have been striving to attain since electrical machinery first came into use.

Jan 2, 1916

By Joseph S. Young

During 1975, the Reclasource Corp. constructed and started up a plant for the recovery of mill scale, coke breeze, blast furnace dust; blast furnace sludge; and TOC dust, at the site of Crucible Steel, Inc., Midland, Pa. The Reclaform Process was employed as the criteria for plant design. This process takes the solid wastes (as listed above) which are commonly found in integrated steel mills, sizes them, forms them through the use of a carbonaceous binder briquetting system, and fixes the binder through a final heat treating step. Blast furnace tests have been performed on 15,000 tons of Reclaform briquets. This paper gives a brief history of the program, and describes the process and the raw materials involved. The blast furnace campaign and its major conclusions are discussed.

Jan 1, 1977

By E. A. Slover

THE usual reverberatory system of smelting cop--1- per concentrate or calcine has for its component parts a furnace and one or two waste heat boilers. These parts are operated on a basis of compromise, since the furnace can send gas to the boilers at too high a temperature and the boilers by plugging, due to dust or slag, can place a definite limit on the amount of fuel the furnace can burn. Over the years the copper concentrate smelting furnace has had few advances in design. The simple rules of design such as the flame should wipe the bath and the speed of the gases should be reasonably low for dust carrying purposes seem to cover the main features. In the construction of the individual furnaces some innovations are always being introduced. Among these are charging so that the work of smelting is a complete bath process, the use of suspended brick arches in place of sprung arches, the use of basic brick, not only in the crucible, but also in roof and sidewalls, the use of various means to feed the charge, the use of magnetite or other heavy material to construct the hearth, water cooling of bridgewall and slag skimming bay, the smelting of raw charge instead of calcine, the use of preheated air, and possibly the use of oxygen-enriched air for combustion. But the general outlines of the furnaces have not changed much except as to size. Furnaces at Hurley As shown on Fig. 1, the furnace at Hurley is 126 ft long between the longitudinal buckstays and 32 ft wide at the skewback plates. The foundation is a concrete retaining wall with piers at intervals that go deeper into the earth. Purposely the wall at the burner end of the furnace is not backed-up as tightly as the other parts of the foundation so that movement due to expansion may take place here rather than into the boiler foundations. Within these foundation retaining walls of concrete, the earth has been removed to allow the placement of the crucible brick base inside of which a silica hearth is laid 4 ft 6 in. in depth. No expansion is left in the brick base and crucible where they are in contact with the hearth. The hearth itself is of quartzite crushed to 1 in. size with fines left in the product. An 8 in. layer is laid and tamped with paving tampers to about 6 in. in thickness. Then a layer of silica flour is spread and vibrated into the hearth. This operation is repeated until a depth of 4 ft 6 in. is occupied by the silica mass onion-skinned in layers of approximately 6 in. Before firing the entire hearth is covered with broken slag to a depth of 4 in. so that a seal may be formed on the hearth. The crucible is completely faced with magnesite chemically bonded brick while the outside, against the foundation, is made of silica brick. The side-walls are carried up with silica brick in which expansion joints are left at intervals. Above the crucible the sidewall is corbelled to form a shelf on which the charge may build up along the side-walls, see Fig. 2. The arch of the furnace is sprung 20 in. silica brick, with the longitudinal centerline horizontal the length of the furnace, and some 9 ft in the center above the bath. Both straight and wedge brick are used in the construction and a thin silica mortar is troweled for joints. After the arch under heat has assumed its permanent shape, a silica slurry is spread over the arch to fill any cracks that have formed, thus giving bearing surface to the brick and preventing dust from entering the body of the arch to act as a future fluxing agent. The uptake of the furnace slopes up to the boiler entrance where a brick pilaster divides the gas stream for the two boilers. Over this flared uptake is a suspended flat arch of firebrick. The pilaster and sidewalls are constructed of firebrick but the bottom of the uptake is lined with silica brick and fettled through holes in the roof with siliceous fettling. Close to the entrance of each boiler is a brick covered slot through which water-cooled dampers may be lowered in event of boiler trouble. These water-cooled dampers are hung permanently in position ready to be lowered when needed. Flexible hoses to follow the dampers as they are lowered are connected at all times and individual chain blocks are used to lower the dampers. A pump supplying water is started before the dampers enter the heat. Charging of the furnace along the sidewalls for some 80 ft from the bridgewall is accomplished by electric vibrating conveyors fed by belt from charge storage bins above the furnace. These conveyors

Jan 1, 1954

By G. C. Rauch, T. H. Courtney, J. Wulff

The superconducting behavior of Nb-Ti alloys containing 40 and higher wt pct of Nb and variable oxygen content was studied as a function of thermomechanical processing. Critical current density (Jc) us applied transverse magnetic field (H) data were obtained for 0.010 in. diam wires at 4.2° Kin steady fields up to 120 kOe. The precipitate responsible for increase in Jc in each of the alloys studied was related to niobium and oxygen content. In alloys containing 40 to 60 wt pct Nb with less than about 1600 ppm of oxygen, w precipitation in the range 350° to 450°C appears to be responsible for the maximum value of Jc observed. In alloys of higher oxygen content (-2500 ppm), the oxygen-enriched a phase which precipitates in the temperature range 450° to 550°C is a more effective fluxoid pinner. Maximum Jc in low oxygen (350 ppm) 80 wt pct Nb alloys was achieved by 600°C aging heat treatments. Measurements of resistive critical field also reflect the changes in composition accompanying precipitation in these alloys. It is well known that the superconducting critical current density, Jc, of solid solution superconducting alloys can be increased by precipitation heat treatments.1-4 The precipitate, to serve as an effective superconzlucting fluxoid pinner and increase Jc, must be of proper size and spacing.' Its electronic properties must also be different from those of the matrix.6 As in age-hardenable systems, optimum temperature and time of aging are dictated by the composition and history of the particular solid-solution alloy under consideration. The alloy diagram of Ti-Nb7-9 shown in Fig. 1 is typical of ß-stabilized Ti alloys. Precipitation to achieve high Jc in Ti-Nb and Ti-Ta alloys is usually carried out in what is believed to be a two-phase a + ß region.1,2,4,10 Previous work on Ti-Nb4,11 and on other ß-stabilized systems such as Ti-V, Ti-Mo,12-14 has shown, however, that the first precipitate to appear is not necessarily (Y. Under specific conditions the meta-stable w phase is precipitated instead. A further complication arises in such alloys from the strong influence of contained oxygen on the phase stability and on the kinetics of precipitation.15,16 The superconducting critical current density of Nb-Ti alloy containing less than 36 wt pct Nb has been studied by Vetrano and Boom1 and Kramer and Rhodes.4 The most effective aging temperature was found to be about 425°C; fluxoid pinning was attributed to the precipita- tion of a by Vetrano and Boom and of w by Kramer and Rhodes. Although the latter workers used recrystal-lized material and the former cold-worked, studies carried out in our laboratory lead us to believe that the type of precipitate responsible for increased Jc is unaffected by prior cold-work.17 That other types of precipitates may also be effective, especially in higher oxygen-containing alloys, has been demonstrated by Reuter, et al.2 in this laboratory with the aid of electron microscopy. Reuter et al. found that at an aging temperature of 600°C, fee-TiO is precipitated, resulting in high values of Jc. Even higher values of Jc were observed after heat treatment at 500°C, but the precipitate could not be indexed as fee-TiO. Whether the precipitate was oxygen-rich w or a could not be clearly distinguished. The work nevertheless indicates that appreciable amounts of oxygen alter the nature of the phases precipitated at different temperatures, at least in alloys containing more than 36 wt pct Nb. Since such compositions are more readily worked into suitable wire than lesser niobium content alloys, it was decided to study the response to aging of the alloy compositions reported in the present paper with the hope of clarifying the effect of composition on the precipitation processes responsible for high values of Jc. I. EXPERIMENTAL Nb-Ti alloys containing 40, 50, 60, and 80 wt pct Nb were prepared by are-melting 35 g ingots of iodide crystal bar titanium and electron-beam melted niobium.

Jan 1, 1969

By W. Klement

By quenching liquid alloys, single-phase solid solutions are obtained in the ranges 0 to 42.0 at. pct Co-Au and 0 to 15 and 75 to about 100 at. pet Co-Cu. Metastable solid solutions are also found in multiphase 42.0 to 49.0 and 69. 0 to about 100 at. pet Co-Au alloys. fie inability to achieve complete solid solubiliiies in Au-co and cu-co alloys is rationalized by a conszderation of the rate processes during cooling and solidification. Trends in lattice distortion and shifts in the compositions of the melting-point minima and critical points are discussed for the binary alloys of copper, silver, and gold with some of the transition metals of the first long period. By rapidly cooling alloys from the melt, a continuous series of metastable solid solutions can often be obtained for components of similar crystal structure which normally form eutectic systems.' The eutectic system Au-Co, Fig. 1, and the peritectic system Cu-Co, Fig. 2, were investigated in order to determine the extent of the metastable solid solutions attainable with the present technique.* EXPERIMENTAL PROCEDURES Most of the alloys were prepared by melting weighed quantities of the elements (purity greater than 99.9 pct) under hydrogen in alumina crucibles by means of induction heating. After reweighing, the alloys were cast under argon into 1 mm diam wires. However, it was not possible to prepare some of the Cu-Co alloys with the requisite homogeneity (over about 1 mm3) in this way. For these alloys, reduced powders (purity greater than 99.9 pct) were mixed for 12 to 48 hr, pressed into compacts that were sintered at 1000º to 1050°C for 6 to 48 hr under hydrogen, and then furnace-cooled. Alloy charges of a few mm3, cut from the cast wires or chipped from the sintered compacts, were heated to 1550º to 1600°C in 30 sec in an alumina insert held within a graphite support and then ejected by a pulse of helium onto a lightly polished copper strip on the inner periphery of a rotating wheel. The

Jan 1, 1963

By R. Rocca, M. B. Bever

THE adiabatic elastic deformation of a body is accompanied by a change in temperature. This phenomenon is known as the thermoelastic effect. Under adiabatic conditions the temperature of a metal bar is decreased by an elastic elongation and is increased by an elastic compression. All materials which elongate on heating behave in this manner. The sign of the thermoelastic temperature change is reversed for the few materials which have a negative coefficient of thermal expansion. The following equation for the thermoelastic effect can be derived from fundamental thermody-namic theorems, as shown in the Appendix: dT8 =-a1/cs T dss [1] Cs In this equation, T is the absolute temperature, S is the entropy, a, is the coefficient of linear thermal expansion, cs' is the heat capacity per unit volume at constant stress and s is the stress (positive in tension). Eq 1 can be applied to small, but finite changes in stress in the following form: ?Ts =- a1/cs T?ss [2] The sign of the temperature change AT is opposite to that of the change in stress s for materials which have a positive coefficient of thermal expansion. Eq 2 has been verified for many different materials and stresses at room temperature. No previous publication appears to have been devoted to an in- vestigation of the thermoelastic effect over a range of temperatures. Review of Literature Lord Kelvin (W. Thomson) presented a general theory of the thermoelastic effect in 1851' and later applied it to a solid, subject to an arbitrary system of stresses'. His analysis was based on a cycle of isothermal and constant-strain transformations and led to Eq 2 above. Even today, one of the clearest expositions of the thermoelastic effect is found in Lord Kelvin's article on elasticity in the Ninth Edition of the Encyclopedia Britannica (1878). R. ROCCA, Junior Member, and M. B. BEVER, Member AIME, are Research Assistant in Metallurgy and Associate Professor of Metallurgy, respectiuely, Massachusetts Institute of Technology, Cambridge, Mass. AIME New York Meeting, Feb. 1950. TP 2760 E. Discussion (2 copies) may be sent to Transactions AIME before APT. 1, 1950, and is scheduled for publication Nov. 1950. Manuscript received Oct. 15, 1949.

Jan 1, 1951

By C. E. Mendenhall

THE word temperature has both a colloquial and a technical use. For everyday purposes of abusing the weather man, no very exact definition is necessary, but for the purpose of giving a simple summary of the physical laws that form the basis of practical pyrometry, something more precise is required. Beginning, therefore, with the common concept of "hotness" and "coldness," we must agree on a method of measuring differences in "hotness," on the unit to be used, and on the point from which measurements are to be taken. We shall then have a definite "scale of temperature," which can be used in all methods of pyrometry. But, as in many similar cases, it is much easier to define or describe something than it is to make practical application and use of the definition; so that much of our attention will be taken up with practical methods of realizing or applying the scale agreed upon. It was early observed that changes in temperature produced large and easily measurable changes in gases, which may be most simply separated into changes in volume (expansion and contraction) under conditions of constant pressure, and changes in pressure under conditions of constant volume. These changes are much the same in magnitude for the common gases oxygen, nitrogen, and hydrogen and also for the rarer helium and argon. This relative uniformity in behavior led to the suggestion of "the gas thermometer" and "the gas scale" as the basis for all temperature measurements. However, as methods were refined, differences appeared between different gases and different ways of using gases, so that Lord Kelvin introduced his more fundamental notion of the "absolute thermodynamic scale" of temperature, which he defined as follows: Given two bodies, say two tanks of water, at different temperatures to determine these temperatures on the "absolute thermodynamic," or Kelvin, scale, operate a thermal engine between these two temperatures, letting it take in heat from the hot body and give out heat to the cold, which therefore corresponds to the boiler and the condenser of a steam engine. The engine we may imagine as a cylinder and piston inclosing a gas and operating with the well-known Carnot or isothermal-

Jan 9, 1919

By A. C. Fieldner, W. A. Selvig

Ash in coal has always becri regarded as an undesirable substance, as the heat content of a coal dccreases in direct proportion to its ash contcnt. It represents so much inert materid that has to be transported, handled when the coal is burned, and finally disposed of. W. R. Roberts, in an address bcfore the Engineers' Society of Western Pennsylvania, Pittsburgh, on Sept. 29, 1925, said that on an annual output of five hundred rnillion tons of coal, a reduction of 3 per cent. of ash means fifteen million tons of refusc and a decrcase of $30,000,000 in freight charges. Ash also often prevents effieient utilization of the coal through the formation of clinker causcd by the mcltirig of the ash constituents when subjected to heat. The mineral constituents of coal may be of such nature as to make the coal unsuited for speeifie uses as for instance, high-sulfur coal for making metallurgical coke. While the usual proximate analysis of coal gives the amount of ash, it gives no information as to the nature and coniposition of the ash nor of the minera1 constituents from which the ash is produced. For the efficient washing of coal to remove ash-forining constituents it is desirable to know the composition and distribution of the ash-fonning constituents. Obviously such information is also desirable for the most efficient utilization of coal. Nature and Composition of Coal Ash Coal ash may be defined as the inorganic residue remailning after cornplete ignitior of coal. It is dcrived from the mineral constituents of thc coal. The ash-forming constituents may be roughly classified as inherent impurities and extraneous impurities. The former are mixed intimately with the coal substance, and arc derived either from the original coal-forming material or from external sources such as sedimentation and preeipitation while the vegetal coal-forming plant remains accumulated. The latter are present in the coal bed as partings, bands or nodules of pyrite! slate, shale, calcit,e, bone, etc. These are formed either during the laying down of the coal bed, or subsequently. Flag-

Jan 1, 1927

By Lawrence Addicks

F. N. FLYNN, Clifton, Ariz.-As a number-of my associates in Arizona know, for a great many years, I have felt that leaching was one of the coming problems. We are about to start experiments at Clifton, along the idea of leaching wet flotation concentrates direct, without roasting-differing in this respect from Mr. Addicks' suggestion. We are prompted in starting these investigations by the difficulty in. transporting and handling flotation concentrates, and consideration for drying and roasting dust losses. I have had the idea for 5 or 6 years past, that those very fine concentrates could best be treated by leaching at the concentrating plant, thus eliminating our transportation difficulties to the smelter, and dust losses at the smelter. Given flotation concentrate, if it must be dried and roasted, it would, in our case, undoubtedly be cheaper to smelt the calcine, but if it can be leached raw and wet and the drying difficulties- and dust loss eliminated, it would be very advantageous to leach it at the mill. I realize that it will be a difficult problem, and one which may take several years to develop, but if developed, it would pay for many years' expense of investigation, and we have been encouraged by our general manager to investigate its merits.

Jan 12, 1916

By C. H. G. Bushell, M. Malnarich

Reagent control in flotation is more an art than a science. Operators vary the amount of re- agents used according to the metallurgy obtained. The amount of collector may be increased, for example, if tailings losses are high and decreased if concentrate grades are low. Often control decisions are based on only one assay, such as the tailings analysis by microscopic or quick assaying techniques. This method of control has given reasonably satisfactory results but is clearly far from perfect. No adjustment in reagent addition can be made until information is received that metallurgy has deteriorated, that is, until the optimum time for adjustment has already passed. The first adjustment probably will be an overcorrection or an undercorrection, and time will be required to establish the best new addition rates. The seriousness of the problem depends on variability of the ore. In ore of continually shifting composition and complexity, mill results will shift back and forth from excessively high concentrate grades and high tailings losses to excessively low concentrate grades and low tailings losses. Economic incentive for improved control also increases with increasing size of mill. Even a small improvement in percentage recovery in a large mill pays for expensive control equipment.

Jul 1, 1956

By Edward H. Robie

IN many ways, the Annual Meeting of the Institute in February was the most successful yet held. Attendance surpassed even that of the 75th Anniversary Meeting in 1947 with its international flavor. The number of technical sessions and the number and quality of papers presented almost certainly set a new record. The banquet was the best attended of any yet held even though the price reached a new high. Despite the demand for almost 1700 tickets to the banquet, some felt that the price-$13.50-was higher than they could afford. There are two points of view: One is that this is the big social event of the year for members and that it is worth from three to five dollars more to have it at the Waldorf and to have a luxury menu, as we did this year, with turtle soup, a fish course, guinea hen for the entree, and wild rice (which we understand is $3.50 per lb) for a vegetable. The other point of view is that a seven or eight dollar banquet of more simple food at the Commodore Hotel would more than make up in appeal, especially to the younger and less well-heeled members, what it would lack in glamour.

Jan 1, 1952

By K. P. Gupta, P. C. Panigrahy

The stabilization of the R, a-Mn, and 0-Mn phases have been studied in the Mn-Ti-Fe and Mn-Ti-Co systems. Iron and cobalt both appear to stabilize the (Mn-Ti) R phase to almost the sarne extent. The R-phase region was found to extend from the lowest e/a to slightly beyond the maximunz e/a limit known for this phase. But, while iron appears to stabilize the a-Mn phase, cobalt tends to stabilize the p-Mn phase. In the two systems manganese appears to get replaced by iron and cobalt in each of the mentioned phases. The instability of the a-Mn phase in the Mn-Ti-Co system and the /3 -Mn phase in the Mn-Ti-Fe system cannot be explained on the basis of adverse size effects because atomic diameters for both iron and cobalt (C.N. 12 at. diam) are ziery similnr and not much different from manganese which they replace. Qualitatively, the reason for the stability of the a-Mn and the p-Mn phases can be traced to the more favorable e/a ratio prevailing in the respective systems and to a competing tendency between the two phases. In transition metal alloy systems the o, p,P, R, a- Mn,' and p-Mn2 phases have been claimed as electron compounds. A large volume of work has been done to establish the criterion for the formation of the o phase but until very recently practically no systematic work was done on the a-Mn and the /3-Mn phases. A recent investigation on the P-Mn phase3 indicates the e/a criterion for p-Mn phase stabilization. Since the R phase was first known to appear only in certain ternary systems1 no detailed work was then possible for this phase. The R phase has been recently discovered as a binary intermetallic compound in the Mn-Ti~ and Mn-si~-' binary systems. The existence of binary R phases opens up the possibilities of studying the effect of alloying elements on the stabilization of the R phase. Of the two binary systems possessing an R phase, the Mn-Ti system appears to be more interesting because at a suitable high temperature it is possible to find the three electron compounds, the a-Mn, p-Mn, and R phases, side by side and it is possible to study the effect of a third transition element on these three electron compounds. For the present investigation iron and cobalt, so called B elements for the formation of electron compounds, have been used as the third element to study the stabilization of the a-Mn, P-Mn, and R phases. EXPERIMENTAL PROCEDURE The alloys were prepared by using 99.9 pct pure electrolytic Fe and Mn, 99.5 pct Co, and crystal bar titanium, supplied by Semi Elements Inc., New York and Gallard Schelsinger Mfg. Co., New York. Weighed amounts of the components were melted in recrystal-lized alumina crucibles in an inert atmosphere (argon) high-frequency induction melting unit. Titanium was made into fine chips for easy dissolution and a special charging procedure was adopted to avoid contacts of titanium chips with the alumina crucibles. Up to 20 at. pct Ti, the maximum titanium content in the investigated alloys, there was no visible sign of reaction of titanium with the alumina crucibles. With a careful control of melting time and temperature the losses were minimized and were always found to be below 0.1 pct. Because of such small and almost constant weight losses, the alloys were not finally analyzed. The alloys were wrapped in molybdenum foil and annealed in evacuated and sealed silica capsules at 1000" * 2°C for 72 hr and subsequently quenched in cold tap water. Annealed samples were examined metallographically and by X-ray diffraction. For all high manganese alloys oxalic acid solutions of various concentrations and 1.0 pct HN03 solution were found suitable as etching reagents. Best contrast between the a-Mn and the R phases could be obtained by using freshly prepared 60 pct glycerine + 20 pct HN03 + 20 pct HF solution. For high iron and cobalt containing alloys, especially for alloys containing the a-Fe, y-Fe, and P-Co phases, 15 cc HNOJ + 60 cc HC1 + 15 cc acetic acid + 15 cc water solution was found to be the best etching reagent. All X-ray diffraction work was carried out (using specimens prepared from annealed powders) with a 114.6 mm diam Debye-Scherrer camera using unfiltered FeK radiation at 25 kv and 10 ma. All calculations for X-ray diffraction work were carried out using an IBM 7044 digital computer RESULTS AND DISCUSSION The two ternary systems, MnTiFe and MnTiCo, were investigated near the manganese rich end, Figs. 1 and 2, and show some common features. In both alloy systems large extensions of narrow R phase regions occur at almost constant titanium contents. At titanium contents higher than that of the single phase R-phase alloys, the same unidentified X phase was found in both ternary systems. The extensions of the X phase close to the Mn-Ti binary indicate that this phase could be the TiMns phase. Too few X phase diffraction lines were present in the diffraction patterns to make positive identification of the X phase. In contrast to this similarity the two systems show opposite behavior in the extensions of the a-Mn and 8-Mn phase regions; while iron tends to stabilize the a-Mn phase, cobalt

Jan 1, 1970

By Philip R. Strand

Vermiculite is the name used for those micaceous minerals with a ferromagnesian aluminum silicate composition and the unique property of exfoliating to a low density material when heated. Commercially, the exfoliated products are also called vermiculite, or more exactly, expanded vermiculite. In its exfoliated state, vermiculite serves many markets, chief among them construction, agriculture, horticulture, and general industry. Composition and Properties Vermiculite, in a natural state, has the characteristic mica habit, splits readily into thin laminae, flexible but inelastic. In a mineralogical sense, vermiculite is a hydrated silicate with no exact chemical composition. The structural formula for a macroscopic, trioctahedral vermiculite may be written: [ ] The structure of vermiculite is basically that of a talc, and it could be regarded as a trioctahedral member of the smectite group. Macroscopically, vermiculite displays the prominent monoclinic crystal faces that are commonly marked by lines at 60° and 120°. Hardness varies from 1.5 to 2.8 or more; specific gravities are between 2.1 and 2.8; color is amber, bronze, brown, dark green, or black. When heated quickly to an elevated tempera¬ture, vermiculite expands by exfoliating at right angles to the cleavage, into wormlike pieces (the name vermiculite is derived from the Latin `vermiculare,' to breed worms). This characteristic of expansion, the basis for commercial use of the mineral, is the result of the mechanical separation of the layers by the rapid conversion of contained water to steam. The increase in bulk volume of commercial grades is 8 to 12 times, but individual flakes may expand as much as 30 times. All expansion occurs at right angles to the basal cleavage. Vermiculite may also be expanded by a number of chemical processes such as soaking in hydrogen peroxide, weak acids, and other electrolytes. There is a color change during expansion that is dependent upon the composition of the vermiculite and furnace atmospheres. Heating in an oxidizing atmosphere produces dull gray colors. The iron is generally not oxidized during chemical expansion. The exfoliation of the vermiculite crystal results in large pores being formed between groups of platelets. Thus exfoliation can make available a large increase in void volume without significantly changing the surface area of the platelets. This characteristic is important in the chemical carrier applications of vermiculite. Although there has been much research on the chemical and structural composition of vermiculite, an exact formula and composition cannot be described. This is to be expected since vermiculite derives from a number of different mineral sources. As has been proposed, vermiculite is not a single mineral species, but families of related minerals. Additionally, commercial interest in the end product has resulted in the name vermiculite being applied to the very voluminous, bloated products from calcination of baueritized (bleached) phlogopite. This has also resulted in many names for the specific vermiculite found in a particular deposit. All of this difficulty arises from the nature of its crystal chemistry. Consequently, vermiculite is now thought of as several related variable mineral series rather than a group of related species. The vermiculites have a common silica tetrahedral sheet structure with separating ions. It is now generally agreed that vermiculites must also be regarded as true clay minerals. The properties of the mineral, such as cation exchange capacity, organic complexing ability, and vari-

Jan 1, 1975

By W. L. Phillips

Stress-strain curves were obtained for single crystals of 70 pct Ag-30 pct Zn tested in tension and shear. Samples tested in tension and shear had comparable resolved shear stresses and stress-strain curves. The {111} <110> slip system was observed. It zoas found that the9.e is a barrier to slip in both latent close -packed directions and that the magnitude of these barriers is proportional to prior strain during easy glide. It was observed that cross-slip in tension and shear was most frequent in crystals with an initial orientation near <100> "Oershoot" zoas observed in tension. The amount of this "overshoot" was independent of initial orientation. AN idealized concept of plastic deformation indicates that a single crystal should yield at some stress that is dependent on crystal perfection and it should then continue to deform plastically by the process of easy glide which is characterized by a linear stress-strain curve and a low coefficient, d/dy, of work hardening. Hexagonal metal crystals generally conform to this ideal concept of laminar flow. In fcc metals the range of easy glide is always restricted in magnitude and it is strongly dependent on orientation, composition, crystal size, shape, surface preparation, and temperature. Since one of the principal differences between the two crystal systems, both of which deform by slip on close packed planes, is the existence of latent slip planes in the fcc crystals, it has been proposed that the transition from easy glide to turbulent flow, characterized by rapid linear hardening, is due to slip on secondary planes intersecting the primary plane.ls Several theories have been proposed to explain the linear hardening and parabolic stages of the stress -strain curve.6"10 The easy-glide region is the least understood of the three stages. The stress-strain characteristics of Cu-Zn, which shows a long easy-glide region, have been extensively investigated."-" In light of recent ideas on dislocations, cross-slip, effect of solute atoms, and stacking fault energy, it was felt that the certain features of this earlier work might be compared with another alloy, Ag-30 pct Zn, which also exhibits a long easy-glide region. Tension and shear stress at room temperatures were employed. The results obtained, together with some interpretation of the observations, are described below. EXPERIMENTAL PROCEDURE The silver and zinc used for mixing the alloys were 99.99 pct pure. The two components were weighed to within 0.1 pct of the weights required fo the alloy composition. They were then placed in a closed graphite mold and the mold and contents were heated in 100°C stages from 500&apos; to 900°C with sufficient time and vigorous agitation at each stage provided to dissolve the silver. The crucible was then heated to 1150°C and agitated violently before being quenched in oil. The resulting alloy rod was machined free of sur face defects and then placed in a graphite mold designed for growing single crystals. The graphite mold was closed with a graphite plug and was encased in a pyrex glass tube which was connected to a vacuum system. The tube and mold assembly were placed in a furnace; the tube was evacuated and the furnace was rapidly heated to a temperature sufficient for fusing and sealing the glass. The glass-encased evacuated mold and contents were then lowered through a vertical furnace. The top section of the furnace was held at 100 °C above the melting point of the alloy. The lowering rate was 1.5 in. per hr. The tension specimens were 1/4 in. diam; the shear specimens were 1/2 in. diam. These specimens were then removed from the mold, etched, and chemically polished with hot (60°C) Chase etch reagent (Crz03-4.0 g, NH4C1-7.5 g, NHOs-150 cc, HzS04-52 cc, and Hz0 to make 1 liter). In preparation for tensile testing, the specimens were carefully machined to a diameter of about 0.200 in. to permit a gage length of 6 in., annealed for 16 hr at 800&apos; to reduce coring, and then cleaned and polished. A modified Bausch-type shear apparatus which has been described previously18 were employed. The gage length was 1/8 in. This shear apparatus was placed in an Instron tensile testing machine. EXPERIMENTAL RESULTS A) Tension. Several specimens were extended at room temperature to determine the effect of initial orientation on the stress-strain curves of Ag-30 pct Zn. The initial orientation and the resolved shear stress supported by the active slip system at various total strains are plotted in Fig. 1. The critical resolved shear stress, t,, initial rate of work hardening, d/dy, and length of the easy-glide region are independent of orientation. The arrival at the symmetry line is shown by an arrow in Fig. 1. During the easy-glide region of the stress-strain

Jan 1, 1963

By F. A. Veer, W. G. Burgers, B. H. Kolster

Starting-- from Cu-Si diffusion couples a study of diffusion in the intermetallic compound Cu3Si of the Cu-Si system has been made. Mmker experiments showed copper to be the diffusing element. From the rate of growth of the Cu3Si phase diffusion coefficients were derived. The activation energy for diffusion is 18.6 kcal per mole. The couples show some interesting microscopic features, for which a tentative explanation is given. THE synthesis of dimethyldichlorosilane, (CH3)2SiC12, an important compound in the manufacture of methyl-silicones, takes place at the surface of a thin Cu3Si layer deposited on a mass of silicon particles."&apos; By passing CH3C1 vapor at about 350°C, silicon atoms are consumed from Cu3Si in forming the volatile silane molecules, thus leaving copper atoms at the surface. It is well-known that copper acts as a poison for the above-mentioned reaction by catalyzing the dissociation of methylchloride.&apos; However, no poisonous effect is observed, at least not before most of the silicon from the particles is converted into silane.3 This would be understandable, if either the copper atoms diffuse through the Cu3Si deposit to the underlying silicon or silicon atoms diffuse in the reverse direction to the copper, both processes resulting in the formation of new Cu3Si and so ensuring the proceeding of the reaction. In this paper experiments are described, demonstrating copper to be the diffusing element. Also a very low activation energy for the diffusion is observed. EXPERIMENTAL The silicon used contained 1 ppm Mg and 1 ppm Cu; the copper was spectroscopically pure (Fe: 2 ppm; Ag, Mg, Si: 1 ppm). The copper and silicon samples were prepared in the form of blocks of approximately 10 by 10 by 5 mm with two parallel sides by spark-machining. The blocks were ground on emery paper (grit 600), etched (copper with nitric acid, silicon with a 4:l nitric acid:hydrofluoric acid mixture), and cleaned with alcohol. As the brittleness of the silicon did not allow a welding procedure to prepare the diffusion couples, the copper and silicon blocks were pressed together in a steel clamp. The pressure exerted by the clamp brought about contact between the parts of the couple and so ensured proceeding of the diffusion. The couples were heated in a vacuum furnace (about 10"5 Torr). The temperature was stabilized within 1 pct. After a heat treatment a section parallel to the direction of diffusion was examined in the microscope and in the X-ray scanning microanalyzer. To determine whether copper or silicon is the diffusing element marker experiments were carried out using tungsten wires with a 20-p diam. EXPERIMENTAL RESULTS Figs. 1, 2, 3, and 4 give micrographs of diffusion zones of Cu/Si couples. Analysis by the electron mi-croprobe of the Cu/Si samples after a definite heat treatment revealed the existence of only one intermetallic phase as can be seen from the line scan shown in Fig. 5. From the X-ray diffraction pattern it could be deduced that this phase was CU 3 Si.4 Even when diffusion zones of considerable width (up to 600 p) were obtained, no other phases occurring in the Cu-Si system could be detected, neither by the electron microprobe nor by X-ray diffraction analysis. The optical anisotropy of the Cu3Si phase made it possible to observe the columnar shape of the crystal-

Jan 1, 1969

By L. D. Jaffe

From data found in the literature, a method has been derived for calculating hardenability of titanium alloys from their composition. A single graph gives the contributions of each alloying element. These are simply added to a base hardenability for unalloyed titanium. Results agree satisfactorily with measured harden-abilities. RECENTLY the importance of hardenability of titanium alloys in permitting attainment of high strength and hardness has become apparent. If a titanium-base alloy has adequate hardenability for the section size and quenching practice used, it can generally be heat treated to high strength and hardness either by simple quenching or cooling from temperatures near the ß transus, by quenching and tempering (aging), or perhaps by direct isothermal transformation of ß1,2 If the hardenability is inadequate, high strength cannot be attained. Preliminary charts translating hardenability requirements for shapes quenched in various media into terms of Jominy end quenched hardenability specimens or ideal round sizes are now available for titanium alloys." The method proposed in 1942 by M. A. Grossmann for calculating hardenability of steels from composition has proved to be of great practical value.&apos; It appeared worthwhile to see whether, using information available in the literature, a corresponding method for calculating hardenability of titanium-base alloys could be derived. Data No systematic experimental study of the hardenability of titanium-base alloys has yet been made. There is available, however, a considerable quantity of scattered data bearing on hardenability. Some investigators determined Jominy curves on one or several alloys. Others reported as-quenched hardness for specimens of a few alloys cooled at various rates or in various media. Still others gave the as-quenched hardness for systematically varied composition, using one or several cooling mediums. The measurements utilized in this work were limited to hardness. Strength measurements were not used because a low strength may reflect high hardness combined with brittleness. Since, in general, the hardness of titanium alloys goes through a peak as the cooling rate is varied from extremely fast to extremely slow, the cooling conditions at which peak hardness was found were taken as the measure of hardenability.1,2 When as-quenched hardness was found to drop as the content of an alloying element (believed to increase hardenability) was raised, this was taken as an indication that the cooling rate was faster than that giving peak hardness; otherwise, the hardness should go up because of solid solution hardening. Also, if the hardness increased on tempering at low temperatures, this was again taken to indicate that the cooling rate had been faster than that giving peak hardness. Many of the data did not indicate precisely the cooling conditions for peak hardness but merely set a maximum or a minimum bound, or a bracket. Since the sizes of specimens quenched and the location of the hardness readings were not always given, it was occasionally necessary to estimate reasonable limits for them. Only the nominal compositions were available in some cases. As quenching from below the ß transus is equivalent to starting with a mixture of a and ß phases, whose individual composition would differ from the overall composition of the alloy, only measurements after quenches from above the ß transus were used. Hardness measurements on samples heated during metal-lographic mounting were also discarded. Because of their bulk, the data used will not be reproduced. They have been or will be published elsewhere.5-22 Approach All data were first translated into terms of ideal round size. For conversion from Jominy or quenched round or sheet, the author&apos;s graphs were used." For a few furnace cools or helium quenches where cooling rates were given, conversion was made on the basis of equal cooling coefficient.23 Plots of ideal round size for peak hardness against percentage of alloying element were next made for

Jan 1, 1956

By K. Tangri, M. Chaturvedi

On water-quenching from within the (a + ß) phase region Zr-2.5 Nb and Zr-2.5 Nb-0.5 Cu alloys can undergo w transfirmation. This transformation has been attributed to the enrichment of ß Zr phase, at the solution-treatment temperature, by the solute atoms. The ß—w transformation is accompanied by a hardening 01 the alloys and in the optimum condition w-bearing specimens of the ternary alloy are about 14 pct harder than those bearing only a&apos; phase. On aging crl reverls thermally to more stable phases, which results in a considerable softening of the specirnens, as follows: w quenched -waged + ß Zr (enriched) waged - a + ß Nb The high-temperature bcc ß Zr in certain zirconium alloys can transform to a metastable w phase which has a primitive hexagonal structure. In the Zr-Nb alloy system w phase has been observed in the composition range of 7 to 15 wt pct Nb by a number of investigators.&apos;-&apos; In this composition range w can form by water-quenching ß Zr or by aging. known as athermal and thermal w, respectively. Hatt and Roberts&apos; have found that thermal w has a constant c/a ratio of 0.622 +0.002 regardless of the composition and tha: the lattice parameters are of the order of a = 5.02A and c = 3.0 ?. The actual lattice parameters depend upon composition and the temperature of aging. The lattice constants of athermal w. however, depend upon alloy composition.3 The presence of w phase in zirconium alloys can cause substantial strengthening which has been attributed as probably due to the coherency between w and parent ß Zr phase. Robinson et al.&apos; found that the presence of w can raise the strength of Zr-5.0 Nb-2.0 Sn and Zr-5.0 Mo-2.0 Sn alloys to about 170.000 psi: however. the ductility is considerably reduced. With the growing use of Zr-2.5 Nb and Zr-2.5 Nb-0.5 Cu alloys in the nuclear reactors, and the significant effect of w transformation on the mechanical properties of zirconium alloys studies were undertaken to study the w transformation in Zr-2.5 Nb and Zr-2.5 Nb-0.5 Cu alloys. This paper is concerned with the formation and reversion of w phase in these alloys. 1) MATERIALS AND EXPERIMENTAL TECHNIQUES The following two alloys were studied: 1) Zr-2.5 Nb-0.5 Cu (referred to as the ternary alloy): 2) Zr-2.5 Nb (referred to as the binary alloy). The chemical analysis of the alloys is given in Table I. The alloys, in the form of 3/8-in.-diam rods and 1/8-in.-thick sheets, were supplied by the Chalk River Nuclear Laboratories of the A.E.C.L. Initial fabrication for preparation of various specimens was carried out by cold rolling and swaging with intermediate anneals at 1000°C. All the heat treatments were carried out after the specimens were wrapped in zirconium foils and encapsulated in silica tubes under a vacuum of 5 x 10-% m of Hg. For optical metallography and hardness measurements disc specimens + in. diam by + in. thick were used. After the required heat treatments the specimens were mechanically and then chemically polished in a 45 pct HNO3, 45 pct H20. and 10 pct HF solution. The hardness measurements on the chemically polished specimens were carried out on a Vickers hardness tester using a 10-kg load. For each specimen at least fifteen indentations were made in order to obtain a representative hardness value. The phase identification and structural analysis were carried out using X-rays and electron diffraction techniques. Wires of 1.5 mm diam reduced to 0.12 mm diam by chemical etching were used for making Debye-Scherrer powder patterns using copper K-a radiations in 114.6-mm-diam camera. Thin films for transmission electron microscopy were prepared by electropolishing heat-treated a by 7 by 0.005-in.-thick strips using a modified Bollmann-Window technique. The 10 pct perchloric acid-90 pct methyl alcohol electropolishing bath was kepl at -5OCC and polishing was done at 5 to 10 v. The thinned specimens were washed in ethyl alcohol at -30" to -40°C and dried between filter papers. The thin films

Jan 1, 1970

R. W. Raymond, New Pork, N. Y.:—I think Mr. Bache has put his finger on the chief source of the danger of dust-, or gas-and-dust, explosions in collieries. 1 mean the persistent determination of the miners&apos; unions to increase their weekly wages by the excessive use of explosives. This would not be feasible if coal-miners were paid by the day; but this form of payment is, for many reasons, not economically practicable; and the universal practice is to pay for the winning of coal according to the quantity produced. If the miner, by using a large amount of powder, can throw down a large amount of coal without corresponding labor on his own part in undercutting and drilling, he will receive more money for less work, provided he is paid for everything—merchantable coal, worth-less dust, slate, and " bone "—resulting from such a method. My attention was called to this matter many years ago by an admirable report of Prof. W. B. Potter, a past- President of the Institute, on the conditions obtaining in this respect in the Illinois coal-field. It was made very clear in that report that considerations of danger to workmen or loyalty to employers could not be relied upon to prevent miners from employing this perilous and wasteful method of increasing their own immediate receipts. So far as I know, only three remedies have been attempted, namely: the enforcement of discipline as to the methods of mining; the refusal to pay for dust, etc., produced by the miners&apos;methods; and restriction upon the use of explosives, effected by requiring the miner to purchase them from the employer, at a price so high as to make it unprofitable for him to use them in excess. All of these attempted remedies have encountered the bitter opposition of the miners&apos; unions. The enforcement of discipline has become almost impossible, if discipline be (as it should be) understood to involve punishment for the violation of rules when no disaster has followed. I

Jan 1, 1910