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By Helmut Thielsch

METALLOGRAPHIC specimens whose surfaces are to be investigated are too thick to allow either light or electrons to pass through them for microexamination by transmission. This difficulty is overcome with optical microscopes by illuminating the surface through the objective lens. For electron microscopes best results are obtained by preparing thin replicas of the surface of the specimen, placing them in the microscope and passing the electron beam through them. Since electron microscopes represent a rather recent invention, they have found little application as compared with optical microscopes. This is true because, aside from the cost of new equipment and development of new techniques, some of the older replica methods necessitated the destruction of the surface from which a thin transparent replica was obtained. In other processes in which the surface is preserved, either too tedious a procedure is required, unsuited to "mass-production " requirements of industrial laboratories, or replicas of insufficient contrast and sharpness are produced. In general, investigations with the electron microscope involve five steps: (I) polishing the surface of the metal specimen, (2) proper etching of the surface, (3) preparing the suitable replica from the surface, (4) examination and photography in electron microscope and (5) interpretation. PREPARATION OF METAL SURFACES Usual metallographic polishing is generally satisfactory. It is important though that the final polish be applied very care- fully, since otherwise evidence of a small amount of deformation might show up even on relatively deeply etched surfaces. A high-power microscope should be used to determine the suitability of the polish and etch. This, of course, is not different from optical examinations, but in these, if a deep etch is used, the problem is less serious. It is frequently difficult to detect with optical microscopes poor polishing aside from superficial scratches before etching. Figs. I and 2 show micrographs of "well" and "poorly" polished nickel etched by immersing the same for 20 min. in a solution containing 8 grams cupricsulphate in 40 C.C. of concentrated HCl and 40 C.C. of H20 (Marble's reagent). The sample shown in Fig. I was polished in its final stages for 5 min. on a carborundum wheel followed by 10 min. on a rouge wheel, whereas the sample shown in Fig. 2 was polished only I min. on the rouge wheel. Generally, this procedure does not guarantee that the surface of the specimen will be either undistorted or deformed. It was applied here merely to show the effects of good and poor polishing. To the experienced investigator, plastic deformation is

Jan 1, 1946

By J. E. Hilliard

A procedure is described for determining, from measurements on a section, the moments of any degree n (i.e., the mean value of Dn, where D is a linear dimension) of the size distribution of particles (or grains) in an opaque sarnple. The analysis is applicable to particles of alnzost any shape and to any mixture of shapes. However, if absolute values of the wloments are required, it is necessary to know what shapes are present and the relative frequency of their occurrence. In addition, certain shape-dependent parameters have to be evaluated before the formulas can be used. The measurement required is the distribution of intercept lengths. This can be determined either visually or with an automatic scanning instrument. Corrections me derived for the bias due to instrumental end effects. The proposed analysis has the following advantages over existing ones: 1) Confidence limits are easily estimated for the measured properties or, conversely, the number of counts required to achieve a given accuracy can be determined. 2) The analysis is not restricted to particles of a special shape. 3) The error resulting from any known degree of departure of the actual from the assumed particle shape can be estimated. 4) The computations are simple and, apart from a change in constants, are independent of particle shape. 5) The number of particles per unit volume can be determined by a simple counting procedure. MANY procedures have been described for determining the size distribution of particles from measurements on a section (or plane of polish). They differ from one another in the property measured on the section (radius, intercept length, profile area) and in the method used for the numerical solution of the integral equation relating the property to the particle size distribution. However, all existing analyses share at least two of the following disadvantages. 1) They are limited to a particular shape of particle (usually spherical). 2) It is very difficult to determine how the statistical sampling error is propagated into the computed size distribution. The investigator thus has no means of knowing the reliability of his result and, what is of equal concern, he is unable to estimate how many measurements are required to obtain a result of the required accuracy. 3) It is also difficult to assess the error resulting from any deviation between the actual and assumed particle shape. The foregoing limitations stem from the complex relationship between the particle size distribution and the distribution of the property measured on the section. As will be shown, they can be avoided by defining the size distribution in terms of discrete parameters

Jan 1, 1969

By Helmut Thielsch

MeTallographic specimens whose surfaces are to be investigated are too thick to allow either light or electrons to pass through them for microexamination by transmission. This difficulty is overcome with optical microscopes by illuminating the surface through the objective lens. For electron microscopes best results are obtained by preparing thin replicas of the surface of the specimen, placing them in the microscope and passing the electron beam through them. Since electron microscopes represent a rather recent invention, they have found little application as compared with optical microscopes. This is true because, aside from the cost of new equipment and development of new techniques, some of the older replica methods necessitated the destruction of the surface from which a thin transparent replica was obtained. In other processes in which the surface is preserved, either too tedious a procedure is required, unsuited to "mass-production" requirements of industrial laboratories, or replicas of insufficient contrast and sharpness are produced. In general, investigations with the electron microscope involve five steps: (I) polishing the surface of the metal specimen, (2) proper etching of the surface, (3) preparing the suitable replica from the surface, (4) examination and phorography in electron microscope and (5) interpretation. Preparation of Metal Surfaces Usual metallographic polishing is generclly satisfactory. It is important though that the final polish be applied very carefully, since otherwise evidence of a small amount of deformation might show up even on relatively deeply etched surfaces. A high-power microscope should be used to determine the suitability of the polish and etch. This, of course, is not different from optical examinations, but in these, if a deep etch is used, the problem is less serious. It is frequently difficult to detect with optical microscopes poor polishing aside fiom superficial scratches before etching. Figs. I and 2 show micrographs of "well" and "poorly" polished nickel etched by immersing the sample for 20 min. in a solution containing 8 grams cupric sulphate in 40 C.C. of concentrated HCl and 40 C.C. of Hz0 (Marble's reagent). The sample shown in Fig. I was polished in its final stages for 5 min. on a carborundum wheel followed by ro min. on a rouge wheel, whereas the sample shown in Fig. 2 was polished only I min. on the rouge wheel. Generally, this procedure does not guarantee that the surface of the specimen will be either undistorted or deformed. It was applied here merely to show the effects Of good and poor polishing. To the experienced investigator, plastic deformation is

Jan 1, 1946

By Helmut Thielsch

MeTallographic specimens whose surfaces are to be investigated are too thick to allow either light or electrons to pass through them for microexamination by transmission. This difficulty is overcome with optical microscopes by illuminating the surface through the objective lens. For electron microscopes best results are obtained by preparing thin replicas of the surface of the specimen, placing them in the microscope and passing the electron beam through them. Since electron microscopes represent a rather recent invention, they have found little application as compared with optical microscopes. This is true because, aside from the cost of new equipment and development of new techniques, some of the older replica methods necessitated the destruction of the surface from which a thin transparent replica was obtained. In other processes in which the surface is preserved, either too tedious a procedure is required, unsuited to "mass-production" requirements of industrial laboratories, or replicas of insufficient contrast and sharpness are produced. In general, investigations with the electron microscope involve five steps: (I) polishing the surface of the metal specimen, (2) proper etching of the surface, (3) preparing the suitable replica from the surface, (4) examination and phorography in electron microscope and (5) interpretation. Preparation of Metal Surfaces Usual metallographic polishing is generclly satisfactory. It is important though that the final polish be applied very carefully, since otherwise evidence of a small amount of deformation might show up even on relatively deeply etched surfaces. A high-power microscope should be used to determine the suitability of the polish and etch. This, of course, is not different from optical examinations, but in these, if a deep etch is used, the problem is less serious. It is frequently difficult to detect with optical microscopes poor polishing aside fiom superficial scratches before etching. Figs. I and 2 show micrographs of "well" and "poorly" polished nickel etched by immersing the sample for 20 min. in a solution containing 8 grams cupric sulphate in 40 C.C. of concentrated HCl and 40 C.C. of Hz0 (Marble's reagent). The sample shown in Fig. I was polished in its final stages for 5 min. on a carborundum wheel followed by ro min. on a rouge wheel, whereas the sample shown in Fig. 2 was polished only I min. on the rouge wheel. Generally, this procedure does not guarantee that the surface of the specimen will be either undistorted or deformed. It was applied here merely to show the effects Of good and poor polishing. To the experienced investigator, plastic deformation is

Jan 1, 1946

By Lucien Eaton

THE accuracy with which detachable bits can be sharpened and gauged by grinding has made possible refinements in design that cannot be attained by forging. This applies particularly to gauging, where accuracy is of the greatest importance. Gauge changes of 1/32 in., and even 1/16 in., are now feasible, whereas 1/16 -in. changes are about the limit in accuracy for forged bits. These small changes in gauge are feasible, however, only with certain fundamental principles of design. Within reasonable limits the depth to which a hole can be drilled is dependent primarily upon the loss in gauge, or diameter, of the bit. In drilling there is always a certain amount of " overdrilling "-the bouncing of the bit in the hole and its rotation wear away the walls-so that the actual diameter of the hole is greater than the diameter of the bit. The amount of this overdrilling depends upon the kind of rock being drilled, and is sometimes sufficient to permit the use of the same diameter of bit for the full depth of the hole. If, for instance, a 1 ¾ -in. diameter bit overdrills 1/32 in. in diameter, and loses 1/64 in. in gauge in drilling 2 ft. of hole, the bottom of the hole is still 1/64 in. larger than the original diameter of the bit, and a new 1 3/4-in. bit can be used to drill the next 2 ft. If, however, the overdrill is less than the loss in gauge, the next bit must be of smaller diameter. The first requisite in reducing loss of gauge in drilling is to have the points and corners of the wings of the bit cut circles of the same diameter, thus producing a reaming surface with continuous bearing on the wall of the hole. The greater the reaming surface, the less the gauge loss that is to be expected. If the reaming surface, however, is too large, it may impede the rotation of the bit through excessive friction. This principle was probably first recognized in the design of the Carr bit (Fig. 1)) a bit with a single cutting edge of wide angle, a definite reaming edge, and a large center hole. The same idea was followed in designing the double-taper cross bits, which are now standard almost everywhere Fig. 2 illustrates a typical double-taper cross bit. The taper from the point of the wing A to a little behind the corner B is 5° and from

Jan 1, 1939

Frontino operates three underground lode mines at Segovia in the area of northwest Antioquia, Colombia, which has been an important gold producing district since the late 1800s. Auriferous pyrite, the pre-dominent sulfide, carries very finely disseminated gold. Other minor sulfides are galena, pyrrhotite, and sphalerite in quartz gangue with the host rock being granodiorite. Silver is associated with the galena and also occurs as a sulfo-salt. In 1.973 and 1974, at an annual milling rate of 145,000 mt (160,000 st), average mill heads were 7.80 g per mt (0.23 oz per st) Au; 20.0 g per mt (0.58 oz per st) Ag; and 0.37% Pb; with a production of 1,150 kg (35,700 oz) Au and 4,500 kg (141,000 oz) Ag. Over 300 mt (330 st) of Pb concentrate were shipped. The ore is crushed to minus 16 mm (5 18 in. ) and fed to a rod mill with the discharge passing over a 1.07 m by 1.07 m (42 in. by 42 in. ) jig operating in closed circuit with the underflow from a cyclone in closed circuit with two ball mills. Cyclone overflow is bulk-floated for sulfide recovery from which the tailings are sent to the river as waste with the flotation concentrate joining the jig concentrate for cyanidation. The combined gold bearing concentrates, after cyclone classification, are reground to a very fine mesh in closed circuit with cyclone classification, with the overflow going to the primary thickener. Regrinding, classification, and thickening is done in cyanide solution, and the primary thickener produces a pregnant solution for Merrill -Crowe precipitation. The thickener underflow is the feed to the agitators, thickener, and filter from which the overflow and filtrate are sent to precipitation. The washed filter cake is repulped and floated for the recovery of a galena concentrate for shipment. About 40% of the recovered silver is in this concentrate. In the refinery the precipitate is melted into a dore bullion. This 600 mt (660 st) capacity mill employs several interesting techniques. Gold recovery of 85% to 90% from a complex sulfide denotes good practice. While total power at 42 kwh per mt (46 kwh per st) is high, it would be recognized that the grinding units are small in size and the flowsheet is complicated due to the gold occurring in the sulfide minerals. The extremely low-cost electric power is generated by

Jan 1, 1981

By Mark Sheppard

DuRing the winter of 1937, the writer visited a West Virginia stone quarry at which the overburden is stripped hydraulically. The quarry is in a bed of limestone, about 200 ft. thick, which outcrops on a hillside above a small river; the quarry face is 175 ft. high. The top of the stone is cut in all directions by erosion channels, which are filled with a tough, red clay. Some of the larger channels are 30 ft. deep and 8 ft. wide. The thickness of the clay overburden ranges from a few inches to 10 ft., the average being estimated as less than 3 feet. The topography of the district is rugged, and the hillside above the quarry face is not suited to the use of the usual mechanical methods of moving dirt. Until hydraulicking was started 10 years ago, the overburden was dug by hand labor, dumped to the quarry floor and transported to the dump in quarry cars. Cleaning deep channels was a slow and expensive process. When hydraulicking was started, the top of the quarry face was over 200 ft. above the level of the river. The pipe line was carried in a semicircle around the top of the quarry about 200 ft. from the face and about 50 ft. above it. As there is no adequate supply of water above the quarry, the water must be pumped from the river. As quarrying progressed the pipe line was moved until the discharge points are now over 300 ft. above the river. Water is pumped from the river by a four-stage centrifugal pump having a 5-in. intake and a 4-in. discharge. It is direct-connected to a 200-hp., 2300-volt, 1760-r.p.m. induction motor. At 1760 r.p.m. the rated capacity of the pump is 750 gal. per minute under a head of 600 ft. The pump is installed in a house of frame construction on the river bank, about 10 ft. above normal water level, which also houses the pump for the mill-water supply. Above the pumps are hoisting bays into which the motors are raised whenever the river reaches flood stage. An 8-in. suction line, equipped with a strainer and a foot valve, takes the water from a concrete sump at the edge of the river. The pipe line from the pump consists of 1750 ft. of 8-in. spiral steel flanged pipe and

Jan 1, 1938

As described in the section on the Palabora twin crushers, the Palabora ore consists of two types: a carbonatite containing a low titanium magnetite, and a foskorite containing high titanium magnetite. As the Palabora pit was widened and deepened, more and more of the foskorite had to be mined. The 25,000 mt expansion of the Palabora plant which went into operation in early 1977 therefore takes its feed from the foskorite section of the orebody. For the expansion, a third 54 in. by 74 in. gyratory crusher was installed to crush feed for two 32 ft by 15 ft (9.75 by 4.6 m), 9000 hp (6710 kw) autogenous single -stage grinding mills. These units are located near the two existing primary crushers, and the ground ore is pumped 5277 ft (1640 m) to the expanded flotation plant by five stages of centrifugal pumps. The new gyratory is installed in a unique Figure-8 shaped enclosure with the crusher in the smaller 27 ft 11 in. (8. 5 m) diameter cylinder and the maintenance area in the 36 ft 1 in. (11 m) diameter cylinder. Cylinder walls are only 23.6 in. (600 mm) thick because their circular design permits all the concrete to be in compression. A rectangular dump pocket which sits on the smaller cylinder contains two manganese steel faced splitters to protect the crusher spider and countersunk water spray headers under the dumping positions for dust suppression. The dump pocket holds 330 st or 300 mt and is also lined with manganese steel. The crusher is mounted on a reinforced concrete diaphragm over the 363 st (400 mt) crushed ore pocket. In the back of the pocket is plastic pipe. Radioactive level control sources are lowered into this pipe from the crusher support floor. The pipe is protected by concrete covered with steel castings. Level control detectors are mounted on the front wall of the pocket. Liner plates on the crushed ore pocket walls are removed and in- stalled by the use of a circular monorail and trolley mounted on the pocket ceiling. The door in the front of the crushed ore pocket is hinged at the top rather than at the sides and is raised by an electric winch. For removal of the eccentric, rails are installed in the pocket and a unique hydraulically operated cart is moved under the eccentric. A platform is raised to receive the eccentric and lowered so that it may be removed. Crushed ore is removed from the pocket by a hydraulically powered, 8 ft

Jan 1, 1978

By Mark Sheppard

DuRing the winter of 1937, the writer visited a West Virginia stone quarry at which the overburden is stripped hydraulically. The quarry is in a bed of limestone, about 200 ft. thick, which outcrops on a hillside above a small river; the quarry face is 175 ft. high. The top of the stone is cut in all directions by erosion channels, which are filled with a tough, red clay. Some of the larger channels are 30 ft. deep and 8 ft. wide. The thickness of the clay overburden ranges from a few inches to 10 ft., the average being estimated as less than 3 feet. The topography of the district is rugged, and the hillside above the quarry face is not suited to the use of the usual mechanical methods of moving dirt. Until hydraulicking was started 10 years ago, the overburden was dug by hand labor, dumped to the quarry floor and transported to the dump in quarry cars. Cleaning deep channels was a slow and expensive process. When hydraulicking was started, the top of the quarry face was over 200 ft. above the level of the river. The pipe line was carried in a semicircle around the top of the quarry about 200 ft. from the face and about 50 ft. above it. As there is no adequate supply of water above the quarry, the water must be pumped from the river. As quarrying progressed the pipe line was moved until the discharge points are now over 300 ft. above the river. Water is pumped from the river by a four-stage centrifugal pump having a 5-in. intake and a 4-in. discharge. It is direct-connected to a 200-hp., 2300-volt, 1760-r.p.m. induction motor. At 1760 r.p.m. the rated capacity of the pump is 750 gal. per minute under a head of 600 ft. The pump is installed in a house of frame construction on the river bank, about 10 ft. above normal water level, which also houses the pump for the mill-water supply. Above the pumps are hoisting bays into which the motors are raised whenever the river reaches flood stage. An 8-in. suction line, equipped with a strainer and a foot valve, takes the water from a concrete sump at the edge of the river. The pipe line from the pump consists of 1750 ft. of 8-in. spiral steel flanged pipe and

Jan 1, 1938

By W. D. Robertson, W. L. Phillips Jr.

Stress-strain curves were obtained for single crystals of alpha brass in tension and in direct shear. Specimens were strained various amounts in a given slip direction, unloaded, and immediately strained in a second slip direction 60°, 120°, or 180&apos; from the original slip direction. Crystals strained in tension and direct shear had comparable critical resolved shear stresses and stress-strain curves. The density of slip lines in direct shear and in tension was essentially the same. The stress-strain curves obtained in shear were independent of initial orientation, choice of {111 } slip plane, choice of <110> slip direction, prior annealing temperature, and rate of cooling after annealing. There was no recovery after annealing for 4 hr at room temperature or 200°C; recovery was observed after 4 hr at 400°C. The crystals showed no asterism and mechanical properties were completely recoverable up to 20 pct strain. It was found that there is a barrier to slip in all latent close-packed directions, and that the magnitude of these barriers, evaluated at 3 pct strain, is proportional to prior strain and independent of the choice of latent direction in the {111} plane. The formation of Cottrell-Lomer barriers is discussed as a possible explanation for the hardening of the latent systems. 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, ds/dE, of work-hardening. Hexagonal metal crystals generally conform to this ideal concept of laminar flow. In face-centered cubic 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 secondary (latent) slip planes in the face-centered cubic 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.&apos;-.; However, the characteristic differences between individual face-centered cubic metals remain to be explained; in particular, it is not clear why the range of easy glide should vary so greatly in different metals and alloys similarly oriented for single slip. An investigation and comparison of different metals with respect to latent hardening on the primary slip plane should provide some of the information required to specify the necessary and sufficient conditions governing the transition from easy glide to turbulent flow. But, in order to accomplish this purpose, plastic strain must be produced by simple shear in a chosen plane and in a predetermined direction by some form of directed shear apparatus, the results of which must be correlated with the corresponding tension experiments. Two such experiments have been performed previously with zinc and with aluminum. Edwards, Washburn, and Parker" and Edwards and Washburn7 found that the strain-hardening coefficients in two latent directions in the basal plane of zinc were the same as in the primary direction. However, to initiate and propagate slip in either the [2110] or the [1210] direction, following primary slip in the [1l20] direction, it was necessary to increase the stress above that required to continue slip in the primary direction; when the direction of shear was reversed 180 deg plastic strain began at a much lower stress than that required to initiate slip in the original direction and the stress to propagate slip in the reverse direction was lower than the stress to continue slip in the forward direction, indicating a permanent loss of strain-hardening. Rohm and Kochendorfer observed softening in aluminum for all latent close-packed planes and directions. They also found that the critical resolved shear stress obtained from their direct shear apparatus was 50 pct lower than the value obtained from conventional tension tests, that the stress-strain curve was linear at 50 pct plastic strain, and that slip lines were not visible at strains less than 30 pct. At present it is uncertain whether these diverse results correspond to real differences in work-hardening characteristics of the close-packed planes of aluminum and zinc or to differences in experimental technique. In view of Read&apos;s analysis &apos;" of the stress distribution in the experimental arrangement of Rohm and Kochendorfer, there is some reason to question the significance of the latter results. In order to resolve this problem it is necessary to re-valuate the direct-shear technique and either repeat the previous measurements or investigate a third system. The latter choice seemed most likely to produce significant results with respect to work-hardening, and accordingly, it was decided to examine the hardening characteristics of the latent slip directions in alpha-brass. The choice of alpha-brass was dictated by the fact that easy glide is more extensive in this alloy than in any other face-centered cubic metal or alloy and, presumably, more nearly like the idealized hexagonal system. Experimental Procedure Crystals were made in graphite by the Bridge-man method in the form of cylinders, 11/2 in. diam and 8 to 9 in. long. Material for the crystals was 70/30 brass containing the following impurities:

Jan 1, 1959

By B. W. Christ, P. M. Giles

Mossbauer effect measurewents have been made on I-mil-thick foils of commercial 1 wt pct C steel and Fe-2 wt pct C alloy. The experimental method required about 3 to 5 vol pct of a phase in the nzultiphase steel sample for detection. Room-temperature Md&apos;ssbauer patterns obtained on austenitized atid quenched samples exhibit fifteen, and possibly twenty-one, lines. A sharp parama&tetic singlet and a quadrupole doublet, poorly resolued from the singlet, are attributed to austenite. Remaining lines are due to tnartensite. Accurate evaluation of austenite line paranzeters is not feasible if significant amounts of other phases such as carbides or martensite occur simultaneously with austenite. This is demonstrated by comparison of hyperfine interactions determined for austenite in multiphase high-carbon samples with those reported for Fe-C austenite in a nearly 100 pct austenitic sanple. Lines from carbides are incompletely resolced from austenite lines, as demonstrated by comparison of austenite line positions with carbide line positions calculated frow published values of hyperfine interactions. One martensite line overlaps an austenite line in the pattern for commercial 1 wt pct C steel. Results of this study suggest that the usefulness of e M6ssbauer pectroscopy for quantitatizle analysis of austenite in bulk samples of quenched and tempered high-carbon steels is restricted by poor resolution. Use of Mossbauer spectroscopy for phase identification and for evaluation of atomic and electronic structures appears quite feasible, however, The Mijssbauer effect has been widely discssed,&apos;-and e Mossbauer effect measurements have been reported on materials of metallurgical interest.7"20 In particular, it has been proposed that e Mossbauer patterns of commercial steels and laboratory-made Fe-C alloys, in the quenched condition, are composed of lines originating in two phases, Fe-C austenite and Fe-C martenite.-&apos; Evidence accumulated in this study demonstrates that three absorption lines found in the central region of the e Mossbauer pattern obtained on quenched steels are attributable to retained austenite. This interpretation is supported by parallel decreases in the intensity of these three lines caused by subambient cooling of commercial 1 wt pct C steel samples after water quenching to room temperature. A second result of this study is to clarify effects of line resolution and sensitivity in the Mossbauer patterns of multiphase steels on the accu- rate determination of austenite line parameters. Experimental line widths (full width at half height) are generally 1.5 to 3 times larger than the natural line width of 0.19 mm per sec. At least two lines, and sometimes more, from a single phase such as cementite (Fe3C), other carbides, martensite, and austenite fall in the energy band, i0.85 mm per sec. hhis band width is employed simply for convenient reference. It represents approximately the energy interval between the + 112 to 112 transitions in ferrite and is expressed as the velocity needed to Doppler shift a 14.4kev 7 ray to the aforementioned ferrite energy levels. This energy band is referred to as "the central region of the Mossbauer atttern:: in this paper. Hence, due to the large number of lines from different phases in a multiphase steel falling in a relatively narrow energy band, absorption lines from the different phases may overlap. Analysis of available data, presented below, indicates that this occurs to a significant extent for phases which commonly occur in quenched and quenched and tempered high-carbon steels. One consequence of limited resolution in the Mdssbauer patterns from multiphase steels is difficulty in accurate determination of such line parameters as position, width, and intensity. In fact, it appears that quantitative analysis for retained austenite in quenched and tempered high-carbon steels is not practical with the present experimental method. Line resolution is influenced to some extent by sensitivity. We point out below that atom or volume fractions of less than about 3 to 5 pct are not detected by the present experimental method. Thus, the presence of a multiplicity of phases does not always lead to impaired resolution. Finally, we report in this paper room-temperature MGssbauer parameters determined for austenite in a freshly quenched, commercial 1 wt pct C steel and in a freshly quenched laboratory heat of an Fe-2 wt pct C alloy. These parameters are compared with others reported in the literature. Three types of hyperfine interactions are detectable in a Mossbauer effect measurement: isomer shift, quadrupole interaction, and magnetic dipole interaction. These interactions are evidenced by one, two, and six line patterns, respectively.&apos;-4 More than one type of interaction has been reported in certain metallurgical phases thus far studied by this method. Isomer shift is the experimentally measured displacement of line position from some arbitrarily defined reference position. In the case of a multiline pattern, isomer shift is given by the displacement of the centroid (center of gravity) of that pattern from the reference position. All isomer shifts measured at finite temperatures contain a second-order Doppler effect characteristic of that temperature. The isomer shift is related to the total s electron density at the nucleus, becoming more negative with increasing s electron density. The first-order quadrupole effect arises from the interaction between the nuclear quadrupole moment and any axially symmetric electric field gradient in

Jan 1, 1969

By H. H. Hausner, H. S. Kalish

M. Hansen—This paper certainly is an interesting study. Although I have not had too much experience in the powder metallurgical methods of studying phase equilibria, I would like to say the following concerning the interpretations of the results obtained: 1. The existence of a zirconium-rich eutectic having a melting point close to 950°C and containing approximately 5 pct beryllium is well established. 2. Undoubtedly sintering of the original compacts (i.e., without repressing and resintering at 1350°C) resulted in a condition being far from equilibrium, even in the low-melting point zircon-rich region where undissolved zirconium particles have been observed. This means that only partial reaction between the component powders has taken place. 3. In preparing and handling powder mixtures for pressing and sintering, we have found that with powders differing considerably in density, and also in particle size, separation in layers of different composition may occur. This means that a concentration gradient would exist within such samples. This phenomenon may, at least to some extent, account for the difference in microstructure of the top and bottom regions of some of the sintered samples. If this is the case, density figures for some of the nominal compositions would not represent actual densities of those mixtures. 4. Fig. 1 shows that the low densities of mixtures with 40 and 60 pct beryllium sintered at 1350°C are changed to much higher densities if the products sintered at 1100°C are repressed and resintered at 1350°C, whereby an approach toward equilibrium takes place. This would mean that the low density and growth in volume is due to nonequilibrium conditions. If this is true, would it be justified, then, to conclude that "the remarkable growth of the alloys in the vicinity of 40 to 60 pct Be indicates the formation of a high-melting point phase, probably accompanied by a considerable change in volume due to a large alteration of the crystal structure from that of the original compounds"? If some compound formation has taken place already during the first sintering at 950" to 1350°C, more compound would be formed by repressing and re-sintering of the 1100" samples. This treatment, however, results in higher, rather than lower, densities. In general, the density-composition curve of alloy systems containing one or more intermediate phases is characterized by a more or less defined contraction (decrease in specific volume, increase in density) over the "theoretical" density. Does not discrepancy exist between the two statements that "growth of the alloy indicates the formation of a high-melting phase . . ." and "even at 1350°C, no indications of sintering have been observed"? 5. I am not sure that the explanation given for the fact that fig. 4 did not reveal as much eutectic as the top portions of the mixture with 2 pct Be, is correct. The density of the melt containing only 5 pct Be or even perhaps less, is not too much different from that of the nominal composition. The reason might be also that there was already some separation of the components in the pressed compact. 6. I do not understand why the microstructure of the bottom regions of the compact with 5 pct Be (fig. 6) is so different from that of the top regions (fig. 5). The compact was melted on sintering at 1100°C. Its composition lies close to the eutectic point. There should be at least some lamellar structure in the bottom regions too; otherwise, the composition of top and bottom must have been very different after sintering, because the eutectic is said to extend as far as the composition ZrBe2. In case the white and gray areas of fig. 6 are both gamma, and the black areas undissolved zirconium, this composition would be close to the phase coexisting with zirconium, that is, ZrBe2, according to the hypothetical diagram, or a compound richer in zirconium. 7. Figs. 9 and 10 are not mentioned in the text. 8. The great difference in microstructure of the composition 20 pct Be of figs. 8 and 14 on one side and fig. 15 on the other side proves that sintering at 950" and 1100°C results only in partial reaction of the powers. 9. The mixture with 60 pct Be (fig. 19) seems to consist of two phases, rather than one phase, one interspersed in a matrix of another. 10. The statement that the eta phase "may be an intermetallic compound or the product of a peritectic or monotectic reaction" seems to be misleading, because the product of a peritectic or monotectic reaction in this region of the system must be an intermetallic compound. 11. If there is some solid solubility of Be in alpha and beta-Zr, it would be expected to be higher in beta-Zr (b.c.c.) than in alpha-Zr (h.c.p.). The temperature of the polymorphic transformation of zirconium then would be lowered, rather than increased. In accordance with this, Battelle has found that the transformation point of titanium is decreased by beryllium. 12. In case the phases present in alloys with 80, 90, and 95 pct Be are identical (which appears to be correct), it is striking that the relative amounts of both phases (eta and beryllium) are not too different within this wide range of composition. With 60 to 65 pct Be

Jan 1, 1951

By J. Walter Snavely

PRESENT-DAY processing in many industries, calcining, sintering, briquetting, beneficiation and nodulizing, increasingly calls for the handling of large volumes of hot bulk materials. Various types of conveyors have been employed. This discussion will cover the factors governing their selection. For temperature ranges up to 400°F, or approximately 200 °C, a wide range of conveyors is available. Special constructions of rubber conveyor belts, steel conveyor belts, vibrating and shaker conveyors, apron conveyors, and drag chain conveyors, all are used successfully. As temperatures go well above 400 2F, however, choice of conveyors is narrowly limited. This paper will consider the problem of handling bulk materials only where the temperatures exceed 400°F. The arbitrary selection of 400 °F as a dividing point undoubtedly can be challenged, as special conveyor belting constructions are available which are suitable for temperatures in excess of 400°F. However, when the relatively short life of such belts and the cost of their replacement, with the attendant down time, are balanced against the reliability and long service life of the properly designed steel constructed units to be discussed, there is little question in any operator&apos;s mind that the special belts are more expensive to use. Because the conveyors under study are for the handling of bulk materials, inevitably including a high proportion of fines, obviously wire mesh belts cannot be included for consideration. Even though this type of conveyor is widely used at high temperatures, i.e., for carrying glassware through a lehr, it is unsuited for the conveying of bulk materials, and therefore will be excluded from further discussion in this paper. Preliminary to the study of the conveyor itself is the determination as to whether the material is to be cooled while it is being handled, or whether the processing requires retention of all heat and the maintenance of a given temperature range. In the majority of cases cooling is incidental to or part of the handling process, when the handling, for example, follows completion of sintering, roasting, calcining, refining, or some other process. To meet such operating conditions successfully, the conveying medium used must have: 1—a construction capable of withstanding maximum initial temperatures of the material being handled. 2—a construction providing efficient heat transfer for cooling. 3—a construction providing dependable operation and long life with minimum service requirements, and 4—a construction providing controlled and efficient conveying. Under the usual conditions of cooling during the handling, the construction selected to withstand the initial maximum temperatures does not necessarily involve using alloys, as excellent results can be achieved with normal carbon steels and cast irons, when they are properly applied and proportioned. The earliest and simplest type of conveyor for handling very hot materials is the cast steel drag chain conveyor, still widely used for handling hot cement clinker, as illustrated by Figs. I and 2. Because of the rugged and generous proportions of the chain link design, low carbon steels are entirely suitable for the links. The pins, however, must be alloy steel. The simple, rugged construction of this type of conveyor makes it readily capable of withstanding high initial temperatures, even though the chain is operating buried in the material. The drag-chain type of conveyor has advantages and limitations. Although the efficiency of the heat transfer is relatively poor, the life of the conveyor is reasonably long, and because of its crude simplicity it does not require much servicing. However, as a conveyor, it is limited in capacity, and largely limited to horizontal runs. Furthermore, because of the crude design, heavy weight, and the chain operating at the temperature of the material, greatly reducing permissible operating chain pulls, this type of conveyor is limited to relatively short centers. Another type of conveyor that has been used for very hot materials is the cast pan conveyor. Because of its very generous proportions the cast pan, which is made of either cast iron or malleable iron, can withstand initial maximum temperatures. It also provides efficient heat transfer for cooling. Further, it is on efficient conveyor construction, which can be used for inclines. Because the chain employs rolling friction instead of sliding friction, and is not in the maximum temperature zone, much longer centers are possible. It is this type of conveyor that is frequently used in the casting of various metal pigs, pig iron, and aluminum; it is obvious, therefore, that very high initial temperatures are being handled. With this kind of conveyor the return run is frequently sprayed with water to accelerate heat transfer. The build-up of residual heat in the very heavy cast pans is thus overcome. The outboard roller steel pan conveyor is an improved pan conveyor&apos; which provides high rates of heat transfer and substitutes formed steel pans for the heavy cast pans. It is a very efficient conveying medium. The details of this particular construction are clearly shown in Fig. 3. An early application of this type of conveyor is shown in Fig. 4. In this case the conveyor units are handling roasted phosphate rock at average temperatures of 1000" to 1500°F, and frequent maximum temperatures as high as 1900°F. Several widths are used. The capacity of the unit at a speed of 50 fpm is approximately 30 tph per inch of width at peak loadings, average capacity being about 1/3 of peak loading. The assembled conveyor is shown in Fig. 5, with views of both the top and the underside to show all the construction details. In particular, the following general design principles were carried out in this construction:

Jan 1, 1954

By B. L. Averbach, M. Cohen

THE PROBLEM MANY hardened steels contain significant quantities of retained austenite even in cases where the carbon and alloy contents are low. In fact austenite has been detected in plain carbon steels containing as little as 0.50 pct carbon.1 Although the austenite may not be visible metallographically it causes definite dilatometric and magnetic effects during subsequent tempering or subatmospheric cooling.1,2,3 Therefore if the reactions which hardened steels undergo during later treatment are to be understood in quantitative fashion it is most essential to have some kind of absolute analysis for the austenite content in hardened steels. Once this analysis has been made, the austenite transformations can be followed precisely by observing such sensitive phenomena as changes in length or in magnetization but the latter techniques can only determine differences in the amount of retained austenite and usually an independent method must be found to put these differences on an absolute basis. In certain instances the microscope can be used advantageously to determine the percentage of retained austenite. The quenched sample is tempered to "darken" the martensite with minimum decomposition of the coexisting austenite, and then is polished and etched in the usual fashion. If the austenite is clearly distinguishable from the tempered martensite and other possible constituents its volume percentage can be measured quantitatively by point counting or lineal analysis.4 Such microscopic methods, however, depend on the etch being sufficiently discriminatory. In fairly coarse structures, if the austenite content is greater than 10-15 pct, there is probably little smearing of the etch, and point counting or lineal analysis seems to be quite accurate. On the other hand if the structure is difficult to resolve, as in commercially treated fine-grained steels which contain less than 10-15 pct retained austenite, the latter may be completely obliterated by the dark etching martensite or it may appear disproportionately low as the interfaces between the two constituents become difficult to distinguish. Another disadvantage of the metallographic method is the need for "darkening" the martensite and hence the determination cannot be applied to as-hardened steels. The required tempering may spoil the structure for other tests or may even remove a small part of the austenite which is being sought. X ray methods hold more promise for covering the range under 15 pct retained austenite. Although austenite and martensite in the hardened steel may be chemically identical, their crystal structures and lattice parameters differ, and the intensity of an austenite diffraction line is some continuous function of the percentage of re-

Jan 1, 1948

By G. C. Gambs

The electric utility industry is the largest customer of the coal industry and consumes nearly 50% of the present coal production in the United States. The byproduct of the burning of coal in utilities is ash in several forms including fly ash, bottom ash and bottom slag. At present, very little of the coal ash produced is being used and must be disposed of as waste. The coal industry should look upon this waste material as a valuable by-product which, if it can be sold for $2 to $4 per ton, will have the net effect of reducing the cost of coal to the utility by 1¢ to 2¢ per million Btu. There are many ways to use coal ash and thus benefit both the coal industry and the electric utility industry by reducing the cost of coal as burned. The solutions to the problems are obvious and all we have to do is apply them. The electric utility industry is the largest customer of the United States coal industry, consuming nearly 50% of present coal production. By 1980, the electric utilities are expected to burn over 500 million tpy of coal or 75% of total coal production. When coal is burned in electric utility plants, ash is produced in one of several forms. These include fly ash, bottom ash and bottom slag. At the present time, approximately 20 million tons of these ashes are produced and must be disposed of by the utilities at a cost of 50¢ to $2 per ton. In 15 years, more than double this quantity will be produced. Only a small quantity of this ash is put to any practical use. Much remains to be done about expanding uses for fly ash and bottom slag. It is well known that fly ash, when used to replace a portion of the cement in a concrete mix, will produce highest quality concrete. The Corps of Engineers and the Bureau of Reclamation have constructed dams and other concrete structures for years with fly ash as a partial replacement for cement. However, many construction firms do not use it. This can become an opportunity for the coal industry to reduce the cost of coal delivered to utilities. Any fly ash, or other ash or slag, which is sold by the utilities, reduces the cost of coal burned by the utility. As much as 1¢ to 2¢ per million Btu could be taken from the fuel cost for each ton of ash sold which otherwise must be disposed of at a cost.

Jan 1, 1967

By M. Cohen, B. L. Averbach

Many hardened steels contain significant quantities of retained austenite even in cases where the carbon and alloy contents are low. In fact austenite has been detected in plain carbon steels containing as little as 0.50 pct carbon.1 Although the austenite may not be visible metallographically it causes definite dilatometric and magnetic effects during subsequent tempering or sub-atmospheric coo1ing.l,2,3 Therefore if the reactions which hardened steels undergo during later treatment are to be understood in quantitative fashion it is most essential to have some kind of absolute analysis for the austenite content in hardened steels. Once this analysis has been made, the austenite transformations can be followed precisely by observing such sensitive phenomena as changes in length or in magnetization but the latter techniques can only determine differences in the amount of retained austenite and usually an independent method must be found to put these differences on an absolute basis. In certain instances the microscope can be used advantageously to determine the percentage of retained austenite. The quenched sample is tempered to "darken" the martensite with minimum decomposition of the coexisting austenite, and then is polished and etched in the usual fashion. If the austenite is clearly distinguishable from the tempered martensite and other possible constituents its volume percentage can be measured quantitatively by point counting or lineal analysis.4 Such microscopic methods, however, depend on the etch being sufficiently discriminatory. In fairly coarse structures, if the austenite content is greater than 10-15 pct, there is probably little smearing of the etch, and point counting or lineal analysis seems to be quite accurate. On the other hand if the structure is difficult to resolve, as in commercially treated fine-grained . steels which contain less than 10-15 pct retained austenite, the latter may be completely obliterated by the dark etching martensite or it may appear disproportionately low as the interfaces between the two constituents become difficult to distinguish. Another disadvantage of the metallographic method is the need for "darkening" the martensite and hence the determination cannot be applied to as-hardened steels. The required tempering may spoil the structure for other tests or may even remove a small part of the austenite which is being sought. X ray methods hold more promise for covering the range under 15 pct retained austenite. Although austenite and martensite in the hardened steel may be chemically identical, their crystal structures and lattice parameters differ, and the intensity of an austenite diffraction line is some continuous function of the percentage of re-

Jan 1, 1949

By B. L. Averbach, M. Cohen

Many hardened steels contain significant quantities of retained austenite even in cases where the carbon and alloy contents are low. In fact austenite has been detected in plain carbon steels containing as little as 0.50 pct carbon.1 Although the austenite may not be visible metallographically it causes definite dilatometric and magnetic effects during subsequent tempering or sub-atmospheric coo1ing.l,2,3 Therefore if the reactions which hardened steels undergo during later treatment are to be understood in quantitative fashion it is most essential to have some kind of absolute analysis for the austenite content in hardened steels. Once this analysis has been made, the austenite transformations can be followed precisely by observing such sensitive phenomena as changes in length or in magnetization but the latter techniques can only determine differences in the amount of retained austenite and usually an independent method must be found to put these differences on an absolute basis. In certain instances the microscope can be used advantageously to determine the percentage of retained austenite. The quenched sample is tempered to "darken" the martensite with minimum decomposition of the coexisting austenite, and then is polished and etched in the usual fashion. If the austenite is clearly distinguishable from the tempered martensite and other possible constituents its volume percentage can be measured quantitatively by point counting or lineal analysis.4 Such microscopic methods, however, depend on the etch being sufficiently discriminatory. In fairly coarse structures, if the austenite content is greater than 10-15 pct, there is probably little smearing of the etch, and point counting or lineal analysis seems to be quite accurate. On the other hand if the structure is difficult to resolve, as in commercially treated fine-grained . steels which contain less than 10-15 pct retained austenite, the latter may be completely obliterated by the dark etching martensite or it may appear disproportionately low as the interfaces between the two constituents become difficult to distinguish. Another disadvantage of the metallographic method is the need for "darkening" the martensite and hence the determination cannot be applied to as-hardened steels. The required tempering may spoil the structure for other tests or may even remove a small part of the austenite which is being sought. X ray methods hold more promise for covering the range under 15 pct retained austenite. Although austenite and martensite in the hardened steel may be chemically identical, their crystal structures and lattice parameters differ, and the intensity of an austenite diffraction line is some continuous function of the percentage of re-

Jan 1, 1949

By George H. Found

DURING recent years considerable interest has been focused on the energy-absorption characteristics of metal when it is cyclically stressed in vibration. The most familiar manifestation of this phenomenon, which is called the "damping capacity" or "internal friction" of a material, is involved in the dissipation of mechanical energy of vibration, which otherwise would build up to stress levels where the problems of noise and fatigue are involved. In addition to its engineering importance in structures, the damping behavior of metals, particularly at low stress amplitudes, has also been used as a unique metallurgical research tool with which to provide sensitive means for indicating certain structural changes within metals. Important contributions toward understanding the energy-dissipating mechanisms that are responsible for internal friction in metals have been made in the last few years. The information is not yet sufficiently extensive, however, to provide a useful metallurgical concept describing exactly how the structure affects the internal friction, or to encourage more general use of this technique among metallurgists. Three fundamental mechanisms have been proposed, as a result of recent experimental investigations, to account for the internal friction in nonferrous metals in low-stress amplitude tests. These include damping due to internal thermal currents that are created and flow in metal because of differential vibratory stress effects,1 damping due to localized inelastic stress relaxation at inhomogeneities within the metal,2,3 and damping related to plastic flow of a localized type for which an explanation has been sought in terms of the dislocation theory.4 How each of these mechanisms functions as a result of structural changes in the metal is, in general, not well understood because of insufficient experimental data for which the controlling variables have been segregated and because of uncertainties in regard to the nature of the structural changes themselves. In several investigations the effects of cold-working on the internal friction of polycrystalline brass and aluminum have been described. As reported by Köster and Rosenthal,5 and by Förster and Breitfield,6 there is an initial large increase in the damping when the metal is worked. The initial high value is not stable, but declines rapidly, at first, in a hyperbolic manner. About 100 hr, after straining, a stable value is reached of magnitude proportional to the amount of strain.7 In Fig. I is shown the relationship found by Zener and his collaborators between the internal friction due to cold-working of alpha brass and the amount of cold-working.8 Several investigators have made internal-friction measurements on worked poly-

Jan 1, 1945

By John R. Lewis

DURING the past few years there has been considerable interest in the sizing and the preparation of the iron ore fed into blast furnaces. Furnacemen know that proper sizing of ore tends to increase the efficiency of the furnace. It has been shown by Joseph1 and his associates that there is a direct relationship between the porosity of an ore particle and the time required for its reduction in the blast furnace. This relationship may be stated as follows: All other factors remaining constant, the time for reduction of a given sample of ore is inversely proportional to its porosity. As a consequence, large pieces of very porous ore, say 3 or 4 in. in cross section, will be completely reduced by the time they reach the hot zone just above the tuyeres, while large pieces of dense, nonporous ore may reach the hot zone with their centers unreduced. This means that for the latter complete reduction will be accomplished by the so-called "direct method;" i.e., where carbon reacts directly with the ore. Some recent work by Philbrook2 seems to contradict this rule. He found that of two unlike ores he studied, the one with the lower porosity reduced more readily. He explained his results as "possibly due to an apparently greater tendency of the `A&apos; ore (lower porosity) to crack during preliminary dehydration and during reduction." This brings his results in line with the rule stated above. It is not the author&apos;s intention to enter into a discussion of the merits of "direct" versus "indirect" reduction. It is enough to say that, in general, indirect reduction is desirable. For the greatest efficiency, Gruner&apos;s theorem3 requires that all reduction in the blast furnace be indirect. Joseph seems to believe this also, because in his recent article on Some Comparisons between Iron Ore, Sinter and Nodules as Blast Furnace Feed&apos; he concludes that: "In general, the carbon in blast-furnace coke is used most efficiently when the largest amount of the carbon charged reaches the tuyeres." However, P. V. Martin,&apos; in an excellent paper, has shown that some direct reduction may actually increase the efficiency of the furnace; that is, increase its rate of production while maintaining maximum fuel economy. For nonporous ores there is more opportunity to have an excess of direct reduction unless the factors that make for indirect reduction are favored; for instance; the proper sizing of the ore particles. Another factor that has an important bearing on the reduction rate, as it is affected by particle size, is the flow of heat into the particle. The ore particle

Jan 1, 1947

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

THE development of high temperature, high stress alloys had proceeded with such rapidity during the war, and for a short time afterward, that our knowledge of the constitution of the alloys had become seriously inadequate. To interpret correctly alloy behavior it is not only necessary to recognize the phases present, but also to know the properties of the phases under all varieties of conditions and treatments. Chromium in these high temperature alloys holds the unique distinction of being the only one element which must be present when an oxidizing or generally corrosive atmosphere is present and when a long life is desired in service. Carbon and nitrogen must also be considered where chromium is present. Carbon is always present, either as an impurity in the raw alloying elements, or as a deliberate addition agent to increase the strength. Nitrogen is generally picked up from the atmosphere during melting or may be added as an alloying element. The effect of nitrogen additions is not too dissimilar to that of carbon additions. As such, chromium carbides and nitrides become of great interest since they are present in a very large number of alloys, including alloy steels, stainless steels, super alloys, and the like. In the chromium-carbon system there are three accepted, carbides: Cr,C, Cr7C8, and Cr3C2. The existence of one more carbide, a still higher carbon form CrC, has been advanced. The phase diagrams of chromium and carbon as published by Friemann and Sauerwald² and by Tofaute, Kuttner and But- tinghaus³ -do not extend beyond the carbide Cr3C2, which contains 13.3 pct carbon; but Hatsuta has extended the diagram up to 20 pct carbon, and in doing so hypothesized, with the use of some questionable experimental results, the compound CrC. This carbide would contain 18.75 pct carbon, 81.25 pct chromium, by weight. It is further known that as the chromium carbides increase in carbon the acid resistance increases but the oxidation resistance decreases4,5 This is of great importance in high temperature alloys where time-temperature stability, as well as oxidation resistance, is of prime importance. The possible existence of a carbide of the formula CrC thus becomes of prime importance and it was considered important to recheck the chromium-carbon binary constitutional diagram for such a compound. In view of the wide spread of values listed for the melting point of chromium (1500° to 1850°C) and 1930°C8, it was considered .that the liquidus temperatures might well be redetermined more ac-

Jan 1, 1951