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By C. Ernest Birchenall

As a basis for further progress in several branches of metallurgy, particularly the study of physical properties of solid solutions and the kinetics of solid-solid reactions, a more complete understanding of the structure of solid solutions and the nature of the metallic bond is necessary. Any investigation contributing to this understanding should be of primary interest to the metallurgist. Since the metallic bond is characteristic of an assemblage of atoms, it can be studied only where this assemblage exists. One of the most promising approaches to this study of the interactions in metals is through the general problem of binary metallic solutions. Although crystal structures of many solid metallic solutions have been determined and solubility limits for many systems are given in phase diagram studies, progress beyond this is irregular. Electromotive force and vapor pressure determinations have provided thermodynamic activity data for small concentration and temperature ranges in the amalgams and a few other systems.1 In two systems, however, copper-zinc and iron-carbon, data are available for the calculation of activities over wide ranges of temperature and concentration. This activity data for the brass system is summarized in. the preceding paper.2 A detailed study of the solubility and distribution of carbon in iron is given by Smith and Darken .3 Their treatment of interaction energies is applicable only to dilute interstitial solutions. The objectives here will be to derive from activity data given in the previous paper for the brasses and structural considerations as much new and detailed information as possible about the interactions between copper and zinc in metallic solid solution and the effect of these interactions in determining the short range structure of the equilibrium phases; that is, we shall investigate the energy of interaction between pairs of atoms of the same and different kinds, the deviations of the distribution of nearest neighbors from complete randomness, and show that the conclusions found are consistent with the observed behavior of the brass system. The method of treatment is fairly general for substitutional metallic solid solutions and will be applied to other systems. The determination of vapor pressures of the components of alloy systems is being undertaken in this laboratory to supply more data for this purpose. The information obtained from solid solutions may be extended by careful analogy to liquids as long as the metallic type of bonding is preserved. Since it will be necessary to employ the language of thermodynamics, a brief summary defining the terms used extensively in solution theory is given below.

Jan 1, 1947

By R. E. Robbins, H. R. Brager

THIS note describes an unusual result obtained in an electron microscope investigation of cavities produced in commercial type 304 stainless steel irradiated in a thermal reactor (ETR) at 290°C to a high neutron flu-ence, 6 x loz1 n per sq cm. The analysis of the steel used is given in Table I. The unusual result is that the calculated quantity of helium gas produced by transmutation during the irradiation, 12 ppm, is less than one-tenth the amount of gas necessary to balance the surface tension of the observed cavities produced in the steel during ETR irradiation. Therefore, the cavities would be expected to behave like voids; however, they behave like "gas filled" bubbles. The maximum quantity of helium produced in the steel by B" transmutation was 8 ppm. In addition, the amount of helium formed by (n, (1) fast neutron reactions with the major alloy constituents and nitrogen was calculated to be 4 ppm based on the tabulated differential or effective fission-averaged cross sections of elements in type 304 stainless steel irradiated in typical thermal reactors.', In previous investigations concerning commercial austenitic stainless steel irradiated in a thermal reactor, the amount of helium gas calculated to be produced by transmutation has always been significantly greater than the amount calculated on the basis that the bubbles were gas filled, PPs= 2y/R. The balance of the gas was assumed to be in bubbles smaller than could be detected on the electron micrographs. Accordingly, the bubbles were attributed to helium produced from transmutation of impurity boron by thermal neutrons. These bubbles were not uniformly dispersed in the matrix but were located principally at grain boundaries, or were radially positioned around a few precipitates located in the matrix and at the grain bundaries., 4 Voids have also been observed in austenitic stainless steel irradiated at high temperature to high neutron fluences in the EBR-I1 and Dounreay Fast Reactors. The voids were uniformly distributed throughout the matrix. The amounts of helium gas produced by transmutation during these irradiations were less than one-tenth the amounts needed to balance the surface tension of the voids formed by the fast reactor irradiation. The fast reactor results are quite analogous to those reported here for the ETR irradiated steel. Electron microscope examination of the fast reactor irradiated material after a 1 hr postirradiation annealing at 900°C disclosed that essentially all the voids had been eliminated. The authors stated that the voids had developed due to the high concentration of vacancies produced during irradiatin., 6, ' The microstructure of the ETR irradiated material as-tested at 75O°C, Fig. 1, contained about 2 X lot5 cavities per cm3 with an average diameter of 110A. With the Van der Waals b correction and y = 1500 ergs per cm2, 300 ppm He would balance the surface tension of the observed cavities. Stereoscopic examination of the electron micrographs showed that the cavities were neither preferentially located near the grain boundaries nor highly concentrated at specific regions or precipitates. The cavities were nearly uniformly distributed throughout the matrix, on dislocations and at grain boundaries. These observations imply that the cavities were not the sole results of B'O

Jan 1, 1969

By H. F. Mills

THE Iron King mine is 12 miles east of Prescott Arizona, in low foothills of the Agua Fria mining district. It was operated in 1906 and 1907, using gravity methods of concentration, mostly on oxide ores. In 1923 and 1924 shipments of sulphide ores were made to the Humboldt smelter. It remained inactive and stripped until 1936, when it was unwatered and sampled to the 400-ft level, the bottom at that time. In 1937 the Iron King Mining Co. purchased the mine and shipped oxide ores. During 1938 a bulk flotation plant of 140 tons daily capacity was erected. The mill has been enlarged and operations have been continuous since the mill first started. In July 1942, the Shattuck Denn Mining Corporation purchased the physical assets from the Iron King Mining Co., and since that time it has operated as the Iron King Branch of the Shattuck Denn Mining Corporation. Economic metals recovered are gold, zinc, silver, lead, a little copper, and fluxing pyrite. The ore occurs as a series of northward-dipping shoots arranged in echelon pattern, the series extending over a length of 3500 ft. Strike lengths of shoots vary from 60 to 600 ft, and widths from 3 ft to as much as 14 ft. The ore occurs in a sheared zone in pre-Cambrian chloritic schist. This schist belt is about 6 miles wide, and the ore zone is about 4 miles from the pre-Cambrian granites to the west. No intrusive rocks have been found in the mine, nor is there evidence of any on surface that might indicate a genetic connection with the ore occurrence (Fig I). The schists strike N23°E and dip 75° to the northwest, which might indicate an overturned thrusting from the granites. In the upper levels, the echelon pattern in which the ore occurs was quite regular, no disturbance of the enclosing schist, or of any faulting that might contribute to the formation of the ore pattern was noted. There is a possibility that the ore zone may be on the axis of an isoclinal fold. Another interesting feature is the more or less typical zoning of minerals along the strike in each of the ore shoots (Fig 2). The north end of each shoot is enclosed by schist, while the south ends usually extend out at full width beyond the economic mining limits in a pyrite-quartz low-grade material. The plane of demarcation between ore and schist is very sharp, with little or no impregnation of sulphide in the schist, especially on the footwall side. The footwall schists are usually firm and strong, and vary from dioritic phases to fine-grained, lighter colored bands, almost silky in texture. The hanging wall for a width of 50 to 8o ft consists of alternating sheetings of chloritic and pyritic schists. The mine is now partly developed to the 92o-ft level, with some diamond drilling from the 1150-ft station of No. 6 shaft. Economically there has been little change in nature or grade of ore from the top of the sulphides, about 200 ft below surface, to the bottom levels. There have been a few echelons in depth, just as there are in plan, and while a few shoots have had

Jan 1, 1947

Secretary's NOTE—on page 176,of this paper, in the fourth line of the first footnote, " 4° " should, be " 1" )'; and on page 180, at the beginning of line 23, ('lined boot" should be " unlined boot." Discussion. Frank Owen, London, Eng. (communication to the Secretary) : In his very interesting and valuable paper, Prof. Snow

Jan 1, 1900

By G. M. Schwartz, J. N. Gundersen

Utilization of magnetite-bearing taconite from the great Mesabi range of Minnesota is fast becoming a major industry in the state. The planning and early development stages of taconite processing projects have been de-scribed1-3 and an excellent discussion has appeared recently4 concerning the current mining, crushing, concentrating, and pelletizing techniques of the Reserve Mining Co., whose operations lie entirely within the Eastern Mesabi district. The present paper will deal only with some of the geological aspects of taconite utilization in the district, with emphasis placed upon some of the common mineralogical and textural features of the tacon-ites currently being mined or stripped and how these factors generally affect the potential recovery of the iron from the metamorphosed Biwabik iron-formation. The generalized geologic setting of the Eastern Mesabi district is presented in Fig. 1. The small circles on the geologic map, adapted from Grout and Broderick,3 mark the location of drill holes that were available for inspection during this study through the courtesy of Reserve Mining and Erie Mining Cos. The cores from these holes were divided into stratigraphic units and examined for their mineral content. Typical mineral associations and grain fabrics were also determined from hundreds of thin and polished sections. Available metallurgical tests made it possible to determine the probable response of a given variety of taconite to magnetic concentration. The longitudinal section of Fig. 1, projected along the strike of the iron formation, presents the simple stratigraphy of the area. It also shows the uniformly thick nature of the iron formation members in the district as well as a somewhat abrupt thickening of the formation to the west of the probable fault. The stratigraphy of these precambrian rocks has been described5-' elsewhere and need only be briefly reviewed. Historically, the oldest rocks unit, the Giants Range granite (Algoman) was covered by the Animikie group of sediments (Middle Huronian?), consisting of the Pokegama, Biwabik, and Virginia formations. Outcrops of the Pokegama quartzite are locally distributed along the northern limit of the Biwabik iron formation, but they are too limited to depict on the geologic map of Fig. 1. The effects of the intrusion of several small diabase sills within the Virginia and Biwabik formations are obscure because all of these rocks have been subsequently highly metamorphosed, mainly by the Duluth gabbro. The emplacement of the transgressing Duluth gabbro (Middle Keweenawan) and that of numerous peg-matitic veins in the iron formation8 has resulted in extensive contact and metasomatic metamorphism which has locally reconstituted the previously existing minerals of the Biwabik formation, in large part quartz and magnetite, into a wide variety of silicate mineral assemblages9 that have a profound effect upon the re coverability of the remaining magnetite. The entire region has subsequently been eroded and largely buried beneath Pleistocene glacial deposits. Structurally, the strike of the gently dipping Animikie group generally parallels the elongate outcrop pattern of the Biwabik iron formation on the map of Fig. 1. The Animikie group most commonly dips from 5" to 15" to the south-southeast. In this area, the Duluth gabbro pluton is somewhat sill-like, dipping about 30" to the southeast. Projecting these present structures upward, it seems probable that the gabbro once extended over the underlying Animikie group in the Eastern Mesabi district. SELECTION AND PRESENTATION OF DATA Because of the locally intense metamorphic activity in the district, mineralogical variations within the magnetite-bearing taconites are commonly abrupt, both vertically and laterally. Consequently it is difficult to select a single hole or small group of holes whose cores will be entirely representative of all of the district. Two holes have been selected, however, that in a general way represent mineralogical and textural variations that are likely to be encountered in both the eastern and western parts of the district. These recently drilled holes were chosen for discussion because it was possible to log the core and then to place the sample intervals at stratigraphic boundaries prior to preparation of the core for magnetic tube tests.

Jan 1, 1961

By Mayumi Someno, Kazuhiro Goto, Masahiro Kawakami

The apparent rates of decarburization of liquid alloys of Fe-C, Fe-C-S, Ni-C, and Co-C systems and the rate of oxidation of solid graphite with pure carbon dioxide gas and with gas mixtures of carbon monoxide and carbon dioxide have been measured in the temperature range of 1000° to 1600°C. The cotnposition of carbon dioxide and carbon monoxide gas at the reaction surface has been measured by oxygen concentration cells with the ZrO,-CaO solid electrolyte. 1) The apparent rates of the carbon removal are essentially the same for all the cases of solid graphite, Fe-C, Fe-C-S, Ni-C, and Co-C systems under the same experimental conditions. 2) The apparent rates are independent of the carbon content in the high carbon concentration range but very much affected by the flow rate and the gas composition of the CO-CO2 reactant gas mixture. The ratio of the gas consumed by the reaction to the total quantity of the supplied gas is very large under the present experi~nental conditions. 3) There is a concentration gradient of' carbon dioxide in the vicinity of the reaction surface and the content of CO, becomes extremely small at the reaction surface. 4) A large time fluc-tuation of the gas composition was observed. This jluctuation implies the presence of unstable flow in the gas phase in the vicinity of the reaction surface. THE decarburization of molten steel by an oxidizing gas or by slag may be one of the most important chemical reactions in steelmaking processes. Nevertheless, the kinetics of this heterogeneous chemical reaction do not seem to be well-solved even with the previous studies. Although the conditions for the reaction in steelmaking processes are quite different from those in the laboratory scale, some critical experiments may give information on the mechanism of the decarburization. From the previous work,'-' it is known that the rate of the decarburization is independent of the carbon content in liquid iron with more than about 0.2 wt pct C when the oxidizing gases are supplied to the surface of liquid Fe-C alloys on a laboratory scale. Two rate-controlling steps have been proposed for the decarburization of liquid iron with the high carbon content: one is the surface reaction control proposed by Swisher and Turkdogan;' the other is that the rate is controlled by the gaseous diffusion through the gaseous stagnant layer. proposed by Baker. Warner, and Jenkins.7 and also by .Ito and Sano.2 In the present study, some experiments have been carried out for the evaluation of these rate-controlling steps in the decarburization of liquid iron with high carbon content. The apparent rate of decarburization of liquid iron has been compared with the rates of carbon removal of liquid Ni-C, Co-C, and solid graphite under the same experimental conditions. The composition of carbon dioxide and carbon monoxide gas at the reaction surface has been measured by oxygen concentration cells. I) EXPERIMENTAL PROCEDURE Fig. 1 shows the schematic diagram of the reaction chamber. Solid graphite and liquid metals were contained in an alumina or magnesia crucible of 32 mm ID and 35 mm in height. The samples were heated by high-frequency induction and the temperature was measured by the calibrated optical pyrometer. The temperature was held constant to within 10°C. The re-actant gases were supplied to the surface of the samples through the quartz tube of 8.0 mm ID. The distance from the end of the quartz tube to the surface was 20 mm. The block of high-purity graphite was cut and shaped to the inner profile of the crucible. The height of the shaped graphite was 18 mm, which corresponded to the depth of the liquid iron of 100 g. About 100 g of Fe-C alloy (4.20 to 4.40 pct C), Ni-C alloy (1.84 pct C), Co-C alloy (1.85 pct C), and Fe-C-s alloy (4.35 pct c, 0.5 or 1.0 pct S) were melted in the crucible. The reactant gases were pure CO, and gas mixtures of CO-CO,: the flow rates were controlled by capillary flowmeters with bleeders.

Jan 1, 1970

By Y. C. Kim, C. B. Manula

This paper is concerned with the details of the derivation of an operations research model, specifically linear programming, to solve production scheduling problems. While some results are presented for an actual study, the calculation of cases of a more general nature have not yet been completed. Production managers from those segments of the mining industry that experience a seasonal or highly variable sales demand find it rather difficult to develop overall producing plans for their organizations. In an attempt to solve this technical problem, demands are usually met by keeping production corresponding exactly to sales. This results in a fluctuating production schedule which is costly to maintain because of overtime premiums in periods of high requirements and because of costs associated with an idle mine plant during slack periods. An alternate solution may be to produce part of the desired amount and make up any deficiencies by using overproduction in previous time periods. This tends to smooth the production pattern through the use of a stockpile. However, because of associated storage costs, the solution may again be undesirable if it yields comparatively large surpluses. In general, this type of scheduling problem has an infinite number of solutions which satisfy the requirements. These are largely dependent on the extent to which an operation is geared to changes in production and on the size of its stockpiling facility. The determination of an efficient schedule is implied, therefore, as one lying between two extreme solutions, i.e., one that minimizes surpluses and one that minimizes output fluctuations. These conflicting objectives result in an economic balance problem between the costs of carrying surpluses forward from slack periods versus the cost of high production levels during peak demands. Various methods of operations research have been developed to handle problems whose genesis is explained above. It is very seldom, however, that these quantitative means are used in practice by the mining industry. Those who are familiar with mining operations know that most managers have not progressed satisfactorily in this area. Production planning systems which most mines employ today are often no more than rough records of management's mental planning. For small mines these means are normally sufficient; but as operations grow larger and become more complex, this type of planning ability is no longer adequate. Managers may find themselves losing control of cost relationships, taking longer to outline production sequences, and becoming forgetful of certain resource availabilities. ORIGIN AND SCOPE OF A SCHEDULING PROBLEM The more fundamental aspects of the production scheduling problem can be brought into focus by demonstrating how it exists in a bituminous coal mine* located in southwestern Pennsylvania. This mine, operating in the Pittsburgh seam, uses a definite room-and-pillar block system of mining, which is planned and carried out with only slight variations to meet local conditions. The coal seam, which is firm and varying in thickness from 8 to 9 ft with the bottom 6% ft being extracted, is under 500 ft of cover and opened by a shaft. The bulk of the coal (60%) is obtained from the pillaring activity with the mains and sections, which are considered primarily as development activities, contributing the balance of 12% and 28%, respectively. All coal is won by a ripper-type continuous miner in conjunction with shuttle-cars directly behind a loading machine. Track haulage is employed to move coal from these machine centers to the shaft bottom. The coal, thus mined, is shipped to one destination as part of a mixture of coals to be used in the making of steel. A material flow diagram of this distribution scheme is illustrated in Fig. 1. PROBLEM STATEMENT From the above types of shipments and mining limitations, the basic problem is to determine an economic means of producing and stockpiling a single coal product produced in three different machine centers (main, section, and pillar). Fig. 2 shows total demand for coal as a function of time. Because of the extremely high peak demands, slack period production must be stored and then distributed

Jan 1, 1969

E. A. Loria (Product Metallurgical Engineer, Crucible Steel Co. of America, Pittsburgh)—In this interesting paper, our introductory work was quoted. We would like to call attention to our sequel paper on the experimental determination of oxygen in cupola-melted cast iron,20 which was not mentioned. Vacuum-fusion oxygen values (as well as hydrogen and nitrogen) were reported for nine heats of cast iron melted in the Battelle 10-in. cupola under normal operating practice and under oxidizing conditions. The oxygen analyses ranged from 12 to 68 pprn compared to the author's computed range of 10 to 80 ppm. The average amount of oxygen found in our irons was about 20 pprn and changes in the silicon content of the iron from 1.32 to 2.35 pct had no consistent effect on the oxygen content of the iron. The gas determination specimens were poured in split steel molds that produced a clean pin, 3/8 in. diam and 2 in. long. Because freezing was almost instantaneous, the pins were entirely white iron (nongraphitic). In the early stages of the investigation, the pins were transferred to a mercury-filled trap system immediately after pouring. This was done to collect gas evolved between pouring and analysis. However, it was found that during storage for 4 weeks gas evolution was negligible. Because the vacuum-fusion analysis was usually completed within 4 days of pouring, pins from later heats were not stored in the mercury-trap system. We found some evidence that cast iron picks up oxygen during long storage, because of rusting. Earlier work by the British Cast Iron Research Association has shown that cast irons may be stored for a long time without significant change in their oxygen content. The practical significance of this study (and our own) would be in the improvement of cast-iron quality. Has the author investigated this aspect and reached any conclusions on the effect of oxygen on the mechanical properties of cast iron? The second phase of our study was to determine the properties of the test bars poured simultaneously with the gas analysis specimens. We realize that there may be complicating factors attendant in this procedure.21 Results from many test specimens measuring chill depth, transverse flexure and deflection strength, spiral fluidity, and sensitivity to hardness of gray irons ranging from 12 to 68 pprn oxygen showed that the lowering of transverse strength was the only significant undesirable effect of high oxygen content. A statistical study of the chill test results21 showed that the iron containing 22 to 46 pprn oxygen had forced chill depths that were 2/32 in. below the expected value from their composition, and irons containing less than 16 ppm oxygen had forced chill depths averaging 1/32 in. greater than the expected chill depth. Higher oxygen contents, within the range of 12 to 68 pprn did not increase forced chill depth. With the wedge tests, there was a good linear relationship between carbon equivalent of the irons and their chill depth. The results indicated that oxygen contents below 50 ppm in the iron did not affect chill depth. With 50 to 70 ppm oxygen in the iron, oxygen appeared to have a slight graphitizing tendency. These results are in disagreement with the common belief in gray iron foundries that "oxidized irons" produce high chill depths. It would be appreciated if the author would comment on this subject. Gustaf Ostberg (author's reply)—In Fig. 1 the legend of line I should read 2 pct C, 1 pct Si. The author wishes to thank Mr. Loria for calling attention to his later work, which was published after the present paper was concluded. The range of oxygen contents quoted seems to agree well with the author's values. The lack of response to variations in silicon content is probably due to the fact that the oxygen content in most cases was below the saturation level. The absence of temperature dependence, even in the case of saturation, is understandable if the difficulty in formation and escape of the deoxidation products is taken into account.

Jan 1, 1960

By J. W. Spretnak, V. Kandarpa

Stress -induced ordering of carbon atoms is studied in a series of Fe-Mn-C alloys. A prominent peak is found in the vicinity of 280°C at frequencies of the order of 1.0 cps, with an associated activation energy of 37 kcal per mole. The height of the peak is linearly rekzted to the concentration of carbon in solution. The distortion of octahedral holes by manganese atoms appears to be predominant over carbon-carbon pair interactions. RELAXATION by stress-induced ordering of point defects is expected whenever the introduction of these point defects produces distortions which have a lower symmetry than that of the lattice. Under zero stress, the isolated point defects occupy the crystallographic-ally equivalent positions in the lattice, as these represent states of equal energy. However, if the defect sites are asymmetric, application of an uniaxial stress will split the energy states, and a redistribution of the defects among various states will take place. This is the case of the internal friction peak called the Snoek peak,1 resulting from isolated interstitials in bcc metals. The interstitial sites in this case have tetragonal symmetry. In the case of fcc and hcp lattices, such an effect is not expected from isolated point defects because of the symmetrical nature of the interstitial sites. However, internal friction peaks arising from interstitial diffusion have been reported both in hcp2,3 and fcc4-8 lattices. These peaks are often explained on the basis of stress-induced ordering of interstitial solutes, caused by the deviation of interstitial sites from their cubic symmetry through the presence of nearby defects. In the case of fcc lattices, evidence for interactions of both the substitutional-interstitial4,6,13 and interstitial-interstitial types5'798'14 have been presented by various investigators. The purpose of the present investigation was to study the internal friction peak attributed to the diffusion of interstitial carbon atoms in high purity austenitic manganese steels and to account for the peak in the light of the existing models. MATERIALS The Fe-Mn-C alloys used in the present investigation were made in two different ways, designated as Type I and Type 11. Type I alloys were made from high purity Fe-Mn alloys obtained in the form of 0.04- in.-diam wires from Materials Research Corporation, Orangeburg, N.Y. These alloys were carburized to different levels using gas mixtures of H2 and CH4 at 1000°C. Type I1 alloys were made in this laboratory starting with zone refined iron, spectrographically pure manganese, and spectrographically pure carbon. They were melted in an argon arc melting furnace and drawn into 0.04-in.-diam wires. All the wires were annealed at about 900°C for 3 hr prior to the internal friction experiments. After the measurements of internal friction, the phases in the samples were identified by X-ray diffraction and the carbon determined by the combustion method. EXPERIMENTAL PROCEDURE In the present work, a classical Ke-type pendulum was used. The details of the equipment were described previously by D. T. Peters.9 Dry helium at 40 torr was used in all the experiments. The internal friction, measured as the logarithmic decrement of the torsion amplitude of vibration was determined as a function of temperature, from ambient to about 500°C. The background internal friction was assumed to have the form of the exponential of the inverse temperature and was subtracted from the raw data. The height of the peak was measured at the position of the maximum in the plot of the internal friction versus temperature. The activation energies of the peaks were measured by the peak shift method. The internal friction values for an alloy were obtained as a function of temperature at different frequencies of vibration. The position of the peak changes with frequency, the higher the frequency the higher the peak temperature. The activation energy of the process associated with the peak is obtained using the formula

Jan 1, 1970

By A. E. Dwight, Paul A. Beck

Known CsCl-type phases in transition metal alloys are shown to be formed preferentially at an electron concentration of approximately 6; a tetragonally distorted version of the CsCl structure is stable at -e/a = 6.5. The relationship of the distortion to electron concentration is verified by experimental data for the quasibinary systems Time ,Os) , Ti(Re,Ir), and Ti(Re,Pt) and the ternary system Ti-Ru-Rh. It was noted some years ago1j2 that the occurrence of CsC1-type ordered structures in equiatomic binary alloys, AB, of transition metals correlates with the location of the component elements in the periodic table. One of the components, A, must be to the left and the other component. B, must be to the right of the chromium group. It was pointed out2 that the CsCl structure does not occur in CrFe. is metastable in VFe,3 and is stable in TiFe up to high temperatures, and that this progressively increasing stability may correspond to an increase in electron transfer between the two components and to a correlated decrease of the magnetic moment associated with the iron atoms. Such a decrease was in fact later demonstrated by Nevitt. 4 The binary CsC1-type phases that are formed by a transition metal with either another transition metal or a lanthanide are listed in Table I. Thirty-three of these phases occur at an average electron concentration (average number of electrons outside of closed shells) of e/a = 6. The five phases that are formed at an electron concentration of e/a = 5.5 appear to be borderline cases. This is suggested for instance by the fact that TiRe is bcc. without any long-range order, while TiTc (at the same electron concentration) has a CsCl ordered structure. Six of the undistorted CsCl phases formed by transition metals with one another have an average electron concentration of e/a = 6. 5. Table I. With the exception of Zr4,Rhs4 and Hf48Rh52. the B component in these phases is a first long-period transition metal, i.e.. iron or cobalt. TiNi. with an electron concentration of e/n = 7, appears to undergo transformation at around room temperature to a more complex ordered struct~re.~ The five other CsCl structures known to occur at e/a = 6.5 have either scandium or a lanthanide metal as the A component; the difference in the elec-tronegativity and in the e/a ratio between the component elements is in these phases very large, A(e/a) - 7, and they may be expected to have a rather strong ionic bond component. Presumably the ionic bond stabilizes the CsCl-type structure at the equiatomic composition and it overrides the electron concentration requirement. When the B element is not from the first long period and the A component is not a rare earth metal. equiatomic (AB) alloys of transition metals with an average electron concentration of e/a = 6.5 in most cases have a crystal structure closely related to the CsC1-type, but slightly distorted tetragonally, Table I. This was first found6 for VRu, NbRu. and TaRu. It was noted6 that the undistorted CsCl structure is also stable in the three systems. but at a lower electron concentration, near e/a = 6. In these systems this requires a deviation from equiatomic stoichiometry. which is otherwise ideally suited for the CsC1-type structure. A similar situation was later found in the Ti-Rh7 and the Ti-Ir systems. 8 The Ti-Rh alloys undergo a second distortion and become orthorhombic when the rhodium content becomes somewhat larger than 50 pct. As mentioned previously. TiRe does not have the CsC1-type ordered structure. In order to investigate how close TiRe is to the borderline of stability of the

Jan 1, 1970

By R. R. Judd, H. W. Paxton

The rate of dissolution of cementite was studied in three low-carbon materials: a zone-refined Fe-C alloy, an Fe-0.5pct Mn-C alloy, and a commercial low-carbon steel. The materials were spheroidized, ad then held isothermally at temperatures above the Al. The isothermal anneal was interrupted periodically by a water quench and the specimens were analyzed by quantitative metallography for the amount of aus-tenite formed during the anneal. The results of this study were compared with an analytical model for the process, which assumes that carbon diffusion in aus-tenite is the rate-controlling step for the cementite dissolution process. The correlation between the model and the experimental data is excellent for the zone-refined Fe-C alloys; however, the Fe-0.5 pct Mn-C alloys and the commercial steel deviate from the calculated model. This deviation is thought to be a result of manganese segregation between the carbide and the matrix. The rate of nucleation of austenite at carbide interfaces was reduced by the manganese addition and enhanced by the presence of ferrite-ferrite grain boundaries. PREVIOUS investigations of the nucleation and growth of austenite from ferrite-carbide aggregates are not entirely satisfying for at least one of several reasons. The most prevalent of these is a lack of quantitative data. Engineering studies have been run on many steels with little control over important parameters such as composition and initial aggregate structure. The data obtained are valid only for material with identical chemistry and thermal history. A more informative approach to the problem of aus-tenitization would be to determine the mechanism that controls the rate of solution of carbide in austenite and how it is modified by alloying elements. This information could then be used to calculate an austeniti-zation rate for any material, provided its composition and structure are known. The object of the present work is to establish the rate-controlling step for cementite dissolution in Fe-C austenite and to investigate the modification of this rate by small manganese additions. The composition and structure of the material used were carefully controlled and all measurements were designed to allow a quantitative analysis of the kinetic process that controls the austenitization rate. A MODEL FOR DISSOLUTION OF CEMENTITE Cementite dissolution has been analyzed mathematically by a model that approximates the material used in the experiments. This model postulates a regular ar-array of identical cementite spheroids with 4 C( diam, embedded in a grain boundary- free ferrite matrix. The analysis provides a detailed description of the dissolution of one carbide spheroid and a generalization of the solution by summation over all the carbides in the material. The carbides may be isolated by defining identical, space-filling cells of ferrite around them. If the cell dimensions are greater than the diameter of the austenite sphere resulting from complete dissolution of the carbide, and no interaction (through diffusion in ferrite) takes place between cells during the dissolution process, the model need concern only one cell, since the solution in each cell is identical. In the experimental material, the dimensions of the cell, the carbide, and the final austenite sphere are approximately 24, 4, and 8 p, respectively; use of the single cell is therefore justified. The experimental observations are made on the austenite nodules that form around each carbide during the dissolution process. The model concerns the growth of these austenite nodules. The attendant shrinking of the carbide can be obtained from the same analysis by an extension of the calculations. Several a priori assumptions are necessary to make the analysis of the growth problem tractable. They are: 1) carbon diffusion through the austenite nodule is the rate-controlling process; 2) local equilibrium exists at all interfaces, 3) the austenite nucleus that forms on each carbide instantaneously envelops the carbide; 4) during the austenite growth process, the diffusion flux of carbon in ferrite is insignificant; 5) a quasi-steady state exists in the austenite concentration field; that is, at any instant during the dissolution process, the austenite carbon concentration gradient closely approximates that for a steady-state solution; and 6) the effects of capillarity on the dissolution rate of the carbides can be neglected. Referring to Fig. 1, a mass balance at the y-a interface for an infinitesimal boundary movement gives: Where rb is the outer radius of the austenite shell, C1 and C are carbon concentrations at the interface in austenite and ferrite, respectively, see Fig. 2, is the diffusion coefficient of carbon in austenite for the concentration of carbon at the interface, and t is time. The fifth assumption permits the austenite carbon concentration to be approximated by the Laplace solution for the spherical case. Therefore, where C(Y) is the carbon concentration at r, and A and B are constants. Local interfacial equilibrium fixes the boundary conditions for the diffusion problem. They are:

Jan 1, 1969

By B. S. Lement

No appreciable stabilization of the austenite-martensite reaction occurs in a 15 pct Cr-0.7 pct C steel unless some martensite is initially present. Stabilization is induced by interrupting the subcool below M,; however, disappearance of the stabilization effect as manifest by burst formation of martensite may occur on continued holding at the final subzero reaction temperature. The higher the temperature of intermediate cycling after the initial subcool, the more permanent is the stabilization effect at the final reaction temperature. Appreciable conditioning of retained austenite occurs on reheating to about 500°C. These phenomena are explainable on the basis of the reaction-path theory. A RECENT investigation' carried out by the authors has shown that isothermal transformation of austenite to martensite definitely takes place in a 15 pct Cr-0.7 pct C steel and that appreciable amounts of martensite can form in this fashion. Investigations2-" arried out on other steels and iron alloys make it seem probable that isothermal formation of martensite is a phenomenon of fairly general occurrence. The reaction-path theory of Cohen, Machlin, and Paranjpe5,7 explains isothermal transformation on the basis of thermal activation of strain embryos which generate martensite plates with a velocity approaching that of sound. On the other hand, the nucleation and growth theory of Fisher, Hollo-mon, and Turnbull8 egards isothermal transformation as due to growth of athermally formed martensite nuclei. The experimental evidence obtained thus far favors the reaction-path theory in that the martensite reaction seems to be nu-cleation-controlled rather than growth-controlled. Stabilization is another phenomenon of importance with respect to the austenite-martensite reaction. According to the reaction-path theory, stabilization is a result of relaxation of strain embryos. It is postulated that strain embryos are present at the austenitizing temperature in the form of dislocations or lattice imperfections in the austenite which vary in the extent to which the lattice is displaced along the strain reaction path leading to martensite. On quenching from the austenitizing temperature, athermal martensite starts to form as soon as the M. is reached, because the activation energy of the transformation which decreases with decreasing temperature is lowered to below the level of the highest energy strain embryos. Martensite plates are generated during the cooling below M, from embryos possessing energies above the activation energy corresponding to the temperature at which the cooling is stopped. The distribution of lower energy embryos is changed because of the elastic and plastic strains produced in the remaining austenite as a result of the transformation. This effectively increases the number of strain embryos at all levels of energy below the activation energy and tends to stimulate further martensite formation either iso-thermally or on subsequent cooling. However, if the temperature at which the cooling is stopped is high enough for relaxation of strain embryos to occur, some degree of stabilization against further transformation may be induced. Whether stabilization will be observed on subsequent cooling will depend on the extent of decrease of activation energy with temperature. If the activation energy decreases sharply with temperature, it is possible that transformation will appear to continue with decrease in temperature in spite of the loss of embryos by relaxation. However, with a more gradual decrease in activation energy, stabilization will be manifest by no further transformation on cooling over a measurable temperature range until the

Jan 1, 1954

By J. T. Norton, N. J. Grant, A. M. Gervais

An investigation of the creep deformation of coarse-grained high-purity aluminum at temperatures approaching the melting point permitted the formulation of a theory of subgrain formation. Subgrain boundaries were found to be formed by two different processes: kinking and polygonization. Polygonization is considered as the main factor in subgrain formation. SINCE "slipless flow" was first discussed in detail by Hanson and Wheeler,' many studies have been undertaken to determine how metals deform at high temperatures. Several types of observations have been reported. Jenkins and Mellor,² working on the creep of mild steels and various grades of commercial irons, found metallographic evidence of a kind of substructure within the grains after deformation at high temperatures. Homes" observed after creep deformation of steel at high temperatures that the original X-ray spots were sharply divided into several small ones. The same phenomenon was reported by Hirst4 from his studies on lead and later by Crussard5 from his work on aluminum. The works of these investigators demonstrated beyond any doubt that the splitting of the X-ray spots was due to the fact that the grains underwent a division into subgrains during deformation at high temperatures. Since then, numerous experiments have been undertaken on aluminum in order to determine the mechanism of formation of subgrains. Research workers can be divided, for the purpose of discussion, into two different groups as regards thinking on subgrain formation. On one hand, based on the work on polygonization by Cahn,6 Guinier and Lacombe,7 Greenough and Smith," Crussard,10,11 and Servi, Norton, and Grant," it was concluded that subgrain formation was a consequence of the simultaneous effects, first, of inhomogeneous deformation which produced bending of the lattice, and second, polygonization. On the other hand, Wood and coworkers13-16 based their work on the results of an earlier report" which dealt with crystallite size and a fragmentation theory. It is interesting to note that while Greenough and Smith8 and Wood et al. performed similar experiments, their conclusions were quite different. Wood et al. did not recognize polygonization as a fundamental mechanism and said that aluminum, at the beginning of the deformation, fragmented immedi- ately into subgrains in order to permit deformation of the metal, the number of these subgrains being a function of strain rate and temperature. They concluded that deformation of grains occurred by flow along the subgrain boundaries. In view of the contradictory conclusions, it would appear that the data were not explicit enough to provide an unambiguous conclusion. It appears that the techniques by which most of the results were obtained, namely, X-ray work, were not enough by themselves to solve the problem. Another approach was needed. In the present research, therefore, an extensive metallographic study was planned to supplement the X-ray work. The importance of revealing the subgrain boundaries metallographically was realized, and this will explain the extensive development and use of etching techniques. In order that subgrains would be considerably larger and therefore more readily visible under a microscope, the creep experiments were performed at very high temperatures approaching the melting point of aluminum. Development of Experimental Technique Creep tests were run at constant load on high purity (99.995 pct) aluminum specimens having a very coarse grain size. None of the specimens was run to fracture. The specimens were originally round (diameter, 0.187 in.; gage length, 1 in.) and had parallel flats milled in the gage section to permit

Jan 1, 1954

By S. Storchheim

MORE and more attention is being paid to the bonding of metals in their solid states. For a better understanding of this technique for joining metals and how it is affected by changes in temperature, pressure, and time at temperature and pressure, a detailed report concerning nickel to aluminum bonding has been published.' In order to broaden the knowledge accrued, some additional work concerning solid state joining of aluminum to copper and aluminum to zirconium was performed. The investigation of the Al-Cu system was considerably more extensive than the investigation of the Al-Zr system. For the A1-Cu system, not only were tensile sudies made but intermetallic penetration rate investigations also were carried out. The effect of temperature on intermetallic penetration rate for the A1-Cu system was determined at 11 tsi pressure, held 2 min. Procedure Apparatus: The hot pressing technique was the means of solid state reaction used and required the equipment depicted in Fig. 1. The following procedure was involved: The two metals to be reacted were placed in an aquadag-lubricated 18-4-1 tool steel die, 16 in. high by 1.440 in. ID, between punches of 1.366 in. diam made of the same material. A thermocouple well was located in the die body 3½ in. down from the top of the die, while another well was located centrally in the bottom punch 8½ in. from the bottom of the die. This die assembly was located in three cylindrical ceramic heating furnaces placed in tandem. Each furnace was controlled individually by a Variac power transformer. In turn, the die and furnaces were placed in a water-cooled stainless steel pot which could be evacuated. A cover, which contained a centrally located Wilson seal with an 18-4-1 1 in. diam ram running through it, was bolted on the pot. After sealing, the pot was evacuated by a roughing pump to 200 microns pressure, after which a diffusion pump was used to bring the pressure down to 5 to 15 u. At this pressure, the furnaces were turned on. As soon as they started to heat, out-gassing of the entire unit raised the pressure to 30 to 400 p. By the time the specimens were at temperature ready to be pressed, approximately 4/2 hr, the vacuum pumps had re-established the 5 to 15 u pressure. Once the desired temperature was reached, the required pressure was applied for a predetermined length of time to the 1 in. ram, through to the top punch, and to the specimen. When the time for keeping the specimen under pressure had elapsed, the pressure was released, the energizing coil current turned off, and the assembly allowed to cool. After cooling, the die was removed from the pot and the specimen was ejected. Specimen Preparation: Two different types of specimens were made for this investigation. One was for subsequent tensile testing, while the other was for determination of intermetallic alloy zone penetration into the parent metals. Tensile Bars—Commercially rolled copper pieces in. thick or zirconium sheet pieces 1/32 in. thick and 1.366 in. diam were placed between commercial 2-S aluminum rod 1 in. thick and 1.366 in. diam. This sandwich in turn was slipped into a 2-S aluminum sleeve 1.438 in. OD and 1.370 in. ID. This sleeve lined the couple up and prevented the aquadag lubricant from getting in between either the A1-Cu or Al-Zr interfaces. Immediately prior to the specimen assembly, the copper or zirconium was abraded on the flat surfaces with 320 grit silicon carbide paper, producing clean smooth surfaces. The aluminum was chemically cleaned just before assembly by: l—degrease in acetone, 2—distilled water rinse, 3—immersion for 3 min in 5 pct NaOH at 70" to 80°C, 4—distilled water rinse, 5—immersion for 2 min in 50 pct HNO3 solution at room temperature, 6—distilled water rinse, and 7—drying in a blast of gas. After the A1-Cu sandwiches were hot pressed and ejected, the specimens were machined such that the aluminum sleeve was removed, and the remaining aluminum was then threaded; the bar so produced was tested later for tensile strength. In all the instances where Al-Cu couples were tested, the specimens broke during the test at the Cu-A1 interface and never within the aluminum or copper. The ultimate tensile strength values at times showed considerable scatter for a set of given reaction conditions. Because of this, as many as three to five specimens were made for a particular set of conditions. The trend of the average tensile strengths obtained was not as conclusive as was the trend of the maximum tensile strengths, the latter values being obtained under optimum reaction conditions. Therefore, the values of ultimate tensile strength given herein are maximums.

Jan 1, 1956

By H. C. Brown

SINCE 1934 the steel industry has been utilizing the spectrograph for supplementing wet chemical analysis in the production control of electric and open hearth furnaces. This means of control made great strides during the war years because of the general acceptance of the spectrograph and the increased emphasis that was placed on rapid control methods. However, in the post war era, with the demand still on increased production, it became apparent that a still more rapid and economical means of production control was needed. Since the spectrograph had been used mostly in the analysis of low alloying and residual elements, it also became apparent that equipment was needed to extend the spectrographic technique to the analysis of the high alloying elements in stainless steel. For these reasons, companies manufacturing spectrographic equipment were prompted to start development work on direct reading instruments. In June 1949, the Applied Research Laboratories of Glendale, Calif., announced that a direct method of spectrochemical analysis for stainless type steels had been developed. This paper will describe the use of the Applied Research Laboratories Production Control Quantometer in the quantitative control of stainless, silicon, and plain carbon steels being made at the Butler Pennsylvania Div. of the Armco Steel Corp. The Armco Butler Div. has one 70-ton electric furnace and six 150-ton open hearth furnaces. The electric furnace is employed in the making of all types of stainless steel and the open hearth furnaces are used for the production of silicon, wheel, and plain carbon steels. The ARL quantometer was purchased primarily for the purpose of controlling the steelmaking in the electric furnace, but its use has been extended for the analysis of final tests (ladle tests) on a number of different types of stainless, silicon, and plain carbon steels. Because of this additional work by the quantometer, substantial savings in manpower and time have been realized by the laboratory. In the analysis of a set of preliminary tests from the stainless steel furnace, approximately 40 min in laboratory time are saved due to quantometric analyses. Despite the fact that more specialty grades of stainless steel are being made in the electric furnace, the average tons per hour have been increased since the quantometer was put into operation. Specialty grades require more furnace time than regular commodity grades of stainless steel. The installation of the ARL production control quantometer was completed on March 13, 1952. By May 1, 1952, the instrument was calibrated for nickel, chromium, manganese, silicon, and molybdenum, which are the elements necessary for the production control of the stainless steel furnace. Within the following month, training of personnel on the quantometer was achieved and a study of the accuracy of the instrument showed that the results obtained were sufficiently accurate for control purposes. Therefore, on June 11, 1952, the quantometer was placed on production control for all types of stainless steels. Starting September 11, 1952, the instrument was gradually placed on ladle analysis (final tests) as the analytical curves were refined and additional curves were drawn. The quantometer has been relatively free of breakdowns since placing it on production control. The samples from only one stainless steel heat have had to be analyzed by wet chemistry because of instrument trouble. The previously existing heat-time record was also bettered by 15 min on a commodity grade of 18-8 stainless steel. Scope of Control In general, the quantometer determines all elements necessary for the production control of all types of stainless steel heats and for the ladle analysis of various types of stainless steel heats. It is also used in reporting final results for silicon, manganese, chromium, nickel, molybdenum, tin, copper, and aluminum on all silicon steel grades and manganese, chromium, nickel, molybdenum, tin, and copper on several plain carbon steel grades. Table I shows the elements and the concentration ranges of these elements in the various types of stainless, silicon, and plain carbon steel that are determined on the quantometer. A study of the results obtained on ladle test samples of stainless steel types 410, 430, 430 Ti, 446, 301, 302, 304, 304L, 305, and 17-7 PH will be discussed. Also included in the study are the results obtained on ladle test samples of a number of silicon steels. Apparatus In order to take full advantage of the potentials of the production control quantometer, the unit has been placed in an air-conditioned room with relative humidity control. The temperature is maintained at 73'22°F and the humidity at 45&5 pct. The air conditioning serves as a precaution to minimize the amount of adjustment and calibration needed during operation. It also reduces contaminating fumes and dust and thereby lessens the necessity for maintenance on the equipment. The quantometer is composed of three units: the high precision multisource unit, the 1.5 meter vertical spectrometer, and the console. The source unit supplies excitation conditions varying from spark-like discharges to arc-like discharges. The voltage to the source unit is supplied by a motor-generator

Jan 1, 1955

By P. R. Sperry, A. P. Beck

It has been shown1 that in poly-crystalline strips of high purity aluminum with a fairly random orientation distribution, grain growth progresses gradually until the average grain diameter reaches a value approximately equal to the strip thickness. Recent work at this laboratory led to the realization that grain growth might be impeded to a considerable extent in the presence of a sharply defined texture, where orientation differences between neighboring grains are small. In order to investigate this effect the following experiment was carried out with the same lot of high purity aluminum previously used for grain growth studies in randomly oriented material.' Very large grain size was developed by grain growth at 650°C in specimens of 0.200 in. thickness. These specimens were then rolled to a thickness of 0.050 in. or 1.25 mm—a reduction of 75 pct. In the rolled strip each large grain corresponded to an elongated area easily identified by etching. After annealing for 1 to 25 min at 600°C and re-etching, these elongated areas were again recognizable. Within each area, corresponding to a single large grain before annealing, there formed by recrystallization a multitude of new grains with a fairly well developed preferred orientation. The orientation and the size of the new grains formed in areas corresponding to different large grains, varied widely depending on the orientation of the parent grains with respect to the rolling direction and the plane of rolling. Many areas were found where the average grain size was considerably smaller than the specimen thickness. Such an area occurred in a specimen cut in half before annealing. One half, containing a portion of the area in question, was annealed 1 min at 600°C, the other half, with the remaining portion of this area, for 25 min at the same tempera-Aluminum killed low carbon steel, § which is now used extensively for severe deep drawing or other difficult forming operations, is unusual in that its grain structure, after cold reduction and box annealing in accordance with conventional continuous sheet or strip mill practice, often is elongated, although at times it is equiaxed. Since this unusual structure has been found superior for many, but not all, severe forming operations, recrystallization of the steel, both at constant temperature and on continuous heating, was investigated and compared with that of rimmed steel in the hope that something might be learned about the mechanism of, and the factors controlling, the formation of such elongated grains. In this structure, the grains are elongated both in the lengthwise direction of the strip and transverse to this direction, even though nearly all of the extension in both hot and cold rolling is in the lengthwise direction. The grains are thus roughly pancake-shaped, being longer and wider than they are thick, as observed also by Burns and McCabe,1 and as illustrated by the typical structures shown in Fig 1. Fig la, representing a conventional longitudinal section, shows the length and thickness of the grains, whereas Fig Ib shows their length and width as seen by examining a section parallel to the sheet surface. Both illustrate the very irregular grain boundaries usually associated with the elongated grain shape. A finer equiaxed grain structure in this same grade is shown in Fig Ic. Either the elongated or the equiaxed structure may be present in the annealed product, and in rare instances the two types may coexist in a single specimen, as shown in Fig 1 d. Isothermal Recrystalliza-tion of Rimmed and Alamimum Killed Steel An aluminum killed steel known to have an elongated grain structure after conventional processing (Steel B, Table l), was selected for the initial recrystallization studies; for comparison, a rimmed steel, A in Table 1, was used. Samples of each in the form of hot rolled strip 0.075 and 0.095 in. thick, respectively, were cold rolled on a small laboratory mill in steps of about 0.010 in. per pass to obtain total reductions of 40 and 60 pct. Small pieces of the cold reduced strip were heated in lead at selected constant temperatures for one of several periods of time, then cooled in air. Rate of heating in the lead was, of course, very fast. Hardness of the cooled specimen was measured and a longitudinal section examined metallographically. Isothermal recrystallization curves for these two steels at 1050°F, based on hardness of the air cooled specimens, are shown in Fig 2 in which the amount of recrystallization corresponding to each plotted point is indicated. The marked difference in the behavior of these two types of steel is evident. After a corresponding amount of cold reduction, the rimmed steel recrys-tallizes in a much shorter time than the killed steel and the shape of its recrystallization curve, (plotted on a logarithmic time scale), is very different. The curve for rimmed steel indicates that recrystallization is analogous to isothermal transformation of aus-i.enite in that it proceeds at a progressively faster rate up to some 50 pct recrystallization, then at an increasingly slower rate. For the aluminum killed steel, however, the start of

Jan 1, 1950

By C. E. Brukl, E. Rudy, St. Windisch

The ternary-alloy system Nb-Mo-C was investigated by means of X-ray, melting point, DTA, and metallo-graphic techniques; a complete phase diagram for temperatures above 1500°C was established. Above 1960°C, niobium monocarbide and the cubic (Bl) high-tenzperature phase in the Mo-C system form an uninterrupted series of solid solutions. The ternary range of the pseudocubic q MoCl-, is very restricted. Dimolyb-denum carbide dissolves up to 44 mol pct Nb2C (2240°C), whereas the maximum solid solubility of MO2C in Nb2C does not exceed 5 mol pct. The order-disorder transformation temperatures in Mo2C and Nb2C are lowered by the mutual metal exchanges. Six invariant (p = const) reactzons occur in the ternary system; three correspond to class 11-type four-phase reactions involving a liquid phase, one to a class I (eutectoid)-type, and two further isotherms are associated with limiting tie lines. The results of the Phase diagram investigation are discussed, and the thermodynamic interpretation identifies the low relative stability of the binary sub-carbides in conjunction with the large stability diflerences between niobium and molybdenum carbides as the cause for the formation of a stable equilibrium between the monocarbide and the metal phase in the ternary reson. Due to their refractoriness, the carbides of the high-melting transition metals have received increased interest in recent years as base materials in composite structures for aerospace applications at high temperatures and for the development of self-bonded cutting tool materials; other novel fields of application include power reactors, where operation at high temperatures becomes essential for attaining high power efficiencies. In these applications, the increased reaction rates at high temperatures require a close consideration of the chemical interactions between the alloy constituents. As a consequence, a detailed knowledge of the phase relationships in the alloy systems is required in order to provide a sound basis for developmental -type work. Partly as a result of the considerable experimental difficulties associated with the investigation of this high-melting alloy class, no complete studies of ternary metal-carbon systems have been performed until recently. Even the high-temperature phase relationships in the binary transition metal-carbon systems have been delineated only during the past few years to a degree of accuracy required for a more detailed study of ternary or higher-order alloys. In recent investigations of binary and ternary systems of refractory transition metals with carbon, boron, and silicon,&apos; alloys from the ternary systems Nb-Mo-C became of interest because of the demonstrated possibility2,3 of obtaining compatible composites based on metal + monocarbide combinations. In the meantime, however, studies in other, but related, ternary metal-carbon systems, such as Ta-W-C, have shown that the solid-state equilibria may change significantly toward higher temperatures (>2000°C), and that extrapolations based on low-temperature equilibrium data are, in general, not very reliable. Although the lower-temperature (<2000°C) phase relationships in the Nb-Mo-C system are similar to those found in Ta-W-C, a cursory thermodynamic analysis of the equilibria indicated4 that complete solid-solution formation between Mo2C and Nb2C should not occur at higher temperatures. The present work was conducted in order to experimentally verify these expectations and, in addition, to provide phase equilibrium data in the melting range of the alloys. In the boundary systems, niobium and molybdenum are known to form a continuous series of solid solutions.576 The continuous solubility was also confirmed by Kornilov and Polyakova,7,8 who also observed a minimum melting point at 22 at. pct Mo and 2345°C. The phase diagram investigations of the Nb-C system by Storms and Krikorian9 and Kimura and Sasaki10 were recently supplemented by Rudy et al.11,12 The system contains a high-melting monocarbide with the B1 structure, Table I, and a subcarbide, Nb2C, which exists in at least two different states of sublattice order at low temperatures47"-&apos;3 and a disordered state above approximately 2500°C.11,12 The melting-point measurements by Rudy et al .11,14 are in close confirmation of the data by Kimura and Sasaki.10 The rather complex phase relationships in the Mo-C system were only recently Clarified.15,18 The system is characterized by three congruently melting, intermediate phases, MozC, ? MoC1-x and a Mol-,, Table I, of which only Mo2C is stable at temperatures below 1650°C. Substoichiometric MozC exists in several states of sublattice order which interconvert in homogeneous phase transitions. Hyperstoichiometric compositions cannot exist in the ordered state. Upon cooling through a critical temperature range, the

Jan 1, 1968

By J. O. Brittain, E. P. Lautenschlager, D. A. Kiewit

Compression tests were run in the temperature range of 700° to 900°C ox 0&apos; phase NiAl intermetal-lic alloys of several grain sizes. At these temperatures the minimum strengths were observed at the stoichiometric composition. While significant increases in strength occurved in both the low-nickel (vacancy) and high-nickel (substitutional) regions, the highest strengths were found in the high-nickel region. During deformation serrated flow was sometimes observed in the low-nickel alloys. After deformation transgranular cvacking and deformation striations were observed in all compositions tested. AS part of a general investigation of the properties of NiAl inter metallic compounds, a preliminary study of the role of point defects upon plasticity was made by high-temperature compression tests on ß&apos; NiAl specimens of several grain sizes and compositions. ß&apos; NiAl is an intermetallic compound having a CsCl structure and a rather wide range of composition from A1-45 at. pct to 60 at. pct Ni.1 According to Bradley and Taylor2 and to cooper,&apos; it possesses a defect lattice in which departures from stoichiometry in the direction of decreased nickel content lead to the presence of vacant nickel sites (although Cooper&apos;s work indicates that a small amount of substitution also occurs) whereas departures on the high-nickel side lead to substitution of nickel on aluminum sites. NiAl forms congru-ently from the melt at approximately 1650°C,1 and thus has a higher melting point than either of its component elements. Up to this time, although this and other high-melting intermetallic compounds have been suggested for elevated-temperature usage,4 only the hardness4 and a few tensile-strength measurements5 have been reported for NiAl at high temperatures. In the present investigation the effects of composition upon the compressive-strength properties in a range of 700° to 900°C have been measured for NiAl of several grain sizes. EXPERIMENTAL PROCEDURES The alloys were made as described elsewhere6 from an A1-46.8 at. pct Ni master alloy furnished by the International Nickel Co. with additions of high-purity nickel and aluminum. The charges were vacuum-induction-melted in A12O3 crucibles with small amounts of helium added to the atmosphere to suppress vaporization. They were cooled slowly from the melting temperature to achieve uniform grain size. In order to refine the as-grown grain size a special rolling technique was developed. Alloys were packed into 0.10-in. wall-type 302 stainless-steel tubes which were partially filled with magnesium oxide to prevent bonding between the alloy and the steel jacket. The ends of the tubes were closed by hot forging, and the packets were then hot-rolled. The alloys with greater than 50 at. pct Ni were rolled at 1100°C, but it was found necessary to increase the temperature to 1350° C before alloys with less than 50 at. pct Ni would roll without cracking. With these temperatures, reductions as high as 48 pct were achieved in a single pass. The rolled alloys will hereafter be referred to as "fine grained" whereas the as-grown material will be designated "coarse-grained&apos;&apos;. The compression specimens were made by cutting square cross-sectional pieces, approximately 3/16 by 3/16 by 1/2 in., with a water-cooled diamond cut-off wheel from the as-grown or the rolled alloys. Specimens were ground to their final dimensions by polishing through 3/0 grit silicon carbide papers. The final shape was a rectangular parallelepiped of square cross section having a height-to-width ratio of 3:1. Compression testing was carried out in a compression rig of our own design mounted on an In-stron Floor Model. The specimen chamber could be heated to 1000°C and was controlled within ±2°C. The compression rig was enclosed within a bell jar and was maintained at a 50 µ of mercury vacuum throughout the duration of the test. The test cham -ber was heated from room to test temperature within 15 min. Specimens were then held at the test temperature 30 min prior to testing. Previous experiments indicated that no grain growth would occur within this time. An Instron Variable Crosshead speed unit was used to adjust for small variations in specimen lengths in order to have a constant initial strain rate, €, for all specimens of a group. For the fine-grained specimens the strain rate was changed rapidly at constant temperature by a factor of 10 with the speed lever on the Instron. For a given € the compression data was analyzed in terms of true plastic strain (E) and true compressive stress (0).

Jan 1, 1965

By A. U. Seybolt

The solubility limit of oxygen in columbium has been determined in the range between 775&apos; and 1100°C by means of lattice parameter measurements and microscopic examination. The solubility is a function of temperature and varies, in the range given above, from 0.25 to 1.0 pct O, respectively. BECAUSE of the marked deleterious effect of oxygen upon the mechanical properties of some of the transition metals, it is desirable to know something about the solubility of oxygen in these metals. The brittleness caused by oxygen in solution is particularly marked in the case of the group VA elements, vanadium, columbium, and tantalum. The solubility of oxygen in vanadium has already been reported in an earlier paper,&apos; and Wasilewski2 has given a value (0.9 wt pct) for the solid solubility of oxygen in tantalum at 1050°C. Brauer3 in 1941 investigated the Cb-0 system up to Cb2O5, but made no real effort to investigate the extent of oxygen solubility in the metal. He made the observation, however, that this solubility must be less than 4.76 atom pct (0.86 wt pct) oxygen. This estimate was made from X-ray diffraction results on the alloys CbO, CbO, and CbO; all alloys consisted of the terminal (Cb) solid solution plus CbO, but the last alloy containing 4.76 atom pct 0 showed only three very weak CbO lines. It is surprising that Brauer, by examining only three alloys, arrived at an estimate of the solubility which agrees very well with the results to be reported herein. Experimental Procedure A columbium strip obtained from Fansteel Metallurgical Products was cut into strips, 0.020x1/2x2 in. Two holes, about 3/16 in. in diameter, were made near the ends of the strips in order to hold them against a flat steel block for mounting in a General Electric X-ray spectrometer for lattice parameter measurements. The same holes were used to hang the specimens inside a fused silica vacuum furnace tube which was part of a Sieverts&apos; gas absorption apparatus. The apparatus and method of adding oxygen gas has been previously described.1 According to the supplier, the columbium obtained had the analysis given in Table I. After degreasing the samples, approximately 0.001 in. was etched from each side of the samples in order to remove possible surface impurities from the last rolling operation. For this purpose the following cold acid pickle was found satisfactory: 8 parts HNO3, 2 parts H2O2 and 1 part HF. Various Cb-O compositions were obtained up to 0.75 wt pct O by the gas absorption and diffusion technique. After the sample had absorbed all the oxygen gas added at 1000°C, an additional 24 hr was allowed for homogenization. This treatment appeared to be adequate, as shown by the linearity of the lattice parameter-composition plot. More concentrated alloys were prepared by arc melting mixtures of Cb and Cb2O5 since it was very time-consuming to make Cb-0 alloys in the neighborhood of 1 pct O, or over, by the diffusion method. When the flat strip specimens were used, they were ready for the X-ray spectrometer after cooling from the Sieverts&apos; apparatus. The cooling rate obtained by merely allowing the hot fused silica furnace tube to radiate to the atmosphere (when the furnace was lowered) was sufficiently fast to keep the dissolved oxygen in solution. Arc-melted alloys were reduced to —200 mesh powder in a diamond mortar, wrapped in tantalum foil, sealed off in evacuated fused silica tubes, and then heat treated as indicated in Table 11. The fused silica tubes were quickly immersed in cold water without breaking the tubes after the heat treatments. The tantalum foil prevented reaction between the fused silica and the sample; there was no reaction between the powdered samples and the foil at 1000°C, but some trouble was experienced at 1100°C. At this temperature level a reaction between the sample and the foil was sometimes observed, which resulted in erroneous parameter values. Experimental Results Hardness Tests: Since most of the X-ray samples were in the form of flat strip, it was convenient to obtain Vickers hardness numbers as a function of oxygen content. Compared to the V-O case,&apos; oxygen hardens columbium much more slowly, presumably because of the larger octahedral volume in colum-bium (about 12.0 compared to 9.3Å3 in vanadium), hence, requiring less lattice strain for solution. The plot of VHN vs wt pct O is shown in Fig. 1.

Jan 1, 1955

By R. R. Estill

THE rank of the Ruhr coal ranges from a high volatile bituminous coal to an anthracite, depending to some extent on the original depth of the seam. The average Ruhr coal corresponds to a soft bituminous American coal of a coking quality. The average thicknesses of individual coal seams being mined are also comparable (59 in. against 65 in. in the United States). However, consideration of seam conditions and mining conditions other than those just mentioned emphasizes differences rather than similarities with United States soft coal. In general, the Ruhr seams now being mined are much more folded and inclined than American seams. Dips of 20&apos; and 30" are common in seams now being worked, and 30 pct of the coal reserves in the district are in seams dipping more than 35". Only on the tops and bottoms of folds do we find rather flat coal seams. In addition to the folding there is extensive displacement by cross faulting plus a certain amount of strike faulting of an overthrust nature, which results locally in doubling or omission of seams. Because of the long history of mining in the Ruhr, nearly all coal lying near the surface has long since been mined out, and we find that the average depth of mining is at present about 2300 ft below the surface. Deep mining, folding, and faulting result in seam conditions requiring a great deal more roof support than one finds in American soft coal mines. In fact only in the anthracite district and the Rocky Mountain and Pacific coal fields do we find somewhat similar conditions. It is easy to say, therefore, that the problem of mechanization of coal cutting and loading in the German mines is quite different from that which we have so effectively met in America with our mobile cutters and loaders, duck bill loaders, and a room and pillar system of mining our drift and slope mines. Partly because of more limited coal reserves, the traditional German mining system is largely the longwall method, which gives an almost complete coal recovery. Backfilling must be extensively practiced to protect the longwall faces, the over and underlying seams and workings, and especially the surface industrialized areas and barge canals. The German engineers have accordingly concentrated their efforts on the design of cutters, loaders, and conveyors suitable to longwall methods rather than room and pillar methods. Undercutters with cutter bars like American models have been in use in the Ruhr since well before World War 11. In 1941 they accounted for 8.5 pct of the production. This percentage, of course, includes coal which was undercut but nevertheless had to be broken down with air hammers or with explosives. The most common of these cutters is the Eickhoff Standard cutter (see fig. 1). This machine does about 95 pct of the undercutting in the Ruhr today, and is available with either compressed air or electrical power and in at least four different sizes. A variation of the cutter is this one with two cutter bars (fig. 2). At the end of 1947 about 200 of these machines and similar cutters were accounting for 13.2 pct of the total production, a production which was, however, only 60 pct of the 1941 production rate, so that the actual cutter tonnage was only up to a small amount over 1941. In 1941 about 3 pct of the production was accounted for by shearing machines making their cut perpendicular to the longwall face. They were similar to those used in the States. These machines are today considered obsolete and now account for only 0.7 pct of the total production. They are located at only a few mines and at present do not seem to have much of a future in the Ruhr. For the future, the Ruhr miner is looking forward to rather extensive mechanization of face work, with two major types of equipment being developed almost simultaneously. On one hand there is the development of cutter loaders for use in relatively hard coal. They represent the further extension of ideas developed after relatively long experience with the Eickhoff cutter. On the other hand there has been since 1942 an intense interest in the Ruhr in the development of face-stripping methods, particularly by the Kohlenhobel (coal plow) and its modification. At the end of 1947 these cutter loaders, Kohlen-hobels and scrapers together were actually accounting for only about 1.4 pct of total production while air hammers still broke 77.1 pct and as much as 1.2 pct was actually broken by hand picks. However,

Jan 1, 1951