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By B. L. Averbach, M. Cohen, C. H. Shih

The isothermal formation of martensite is studied in Fe-Ni-Mn and Fe-Mn-C alloys under conditions where the athermal transformation is completely avoided, there being no martensite present at the beginning of the isothermal reaction. The initial rate of martensite formation is extremely low under these conditions, but it increases greatly as soon as some martensite is formed. It appears that true initial characteristics of the isothermal reaction can be studied only in samples wherein the formation of martensite on cooling is completely eliminated. THE phenomenon of isothermal martensite formation is well established1 -" but the observations have been made in the presence of varying quantities of athermal martensite. Complete avoidance of athermal martensite and the formation of martensite under pure isothermal conditions have been reported by Kurdjumov and Maksimova2 in an Fe-Ni-Mn alloy (23 pct Ni, 3.4 pct Mn, and the balance Fe) and have been confirmed by recent experiments of Cech and Hollomon in a similar alloy. In each case, the isothermal transformation rate was a maximum at the beginning of the transformation and decreased as the reaction proceeded. These initial transformation rates were used by Fisher7 for the confirmation of calculated nucleation rates based on classical nucleation theory extended to include the effect of elastic strain energy. Cech and Hollomon have admitted, however, that up to 3 pct martensite formed on cooling in their samples and thus was present at the initiation of their isothermal runs. In this paper, experiments on similar alloys are described, except that care was taken to remove the incidental martensite developed at the surface prior to the isothermal reaction. It is shown here that the reaction rate at the beginning of the strictly isothermal reaction is very low. As soon as some detectable amount of martensite is produced, the transformation rate increases rapidly but ultimately decreases as the reaction proceeds. When martensite is allowed to form isothermally in the presence of about 2 pct of prior martensite, the apparent initial reaction rate is a maximum and corresponds to the observations of Kurdjumov and Maksimova and Cech and Hollomon. It is evident that the cooperative nature of the martensitic reaction results in a significant stimulating or autocatalytic effect on the isothermal reaction rate. Experimental Procedure The chemical compositions of the alloys used in this investigation are listed in Table I. The Fe-Ni-Mn alloys A and B were melted in an induction furnace, cast into ingots 2 in. diam by 4 in. long, and forged into % in. diam rods. The manganese steel (alloy C) was melted in an arc furnace and cast into ceramic molds 1/4 in. diam by 5 in. long. The rods were sealed in evacuated Vycor tubes and homogenized at 1175" to 1200°C for 50 hr. The long homogenization resulted in some manganese loss from the surface, producing a thin layer of transformed material. This surface layer was removed by grinding and the specimens were then swaged to 0.050 in. diam wire. Measurements of electrical resistance were made with a Kelvin double bridge to follow the course of the martensitic transformation. A small permanent magnet was also useful in scanning an entire specimen for small localized quantities of martensite. This test is almost as sensitive as the electrical resistance method in detecting the presence of martensite (which is ferromagnetic in these alloys) and it was especially useful in deciding whether a specimen was completely free of martensite. The wire specimens (0.050 in. diam by 2.5 in. long) were sealed in evacuated tubes and austenitized for 5 min at 1150°C. The specimens were quenched into water and, unless otherwise stated, stress-relieved for 1 hr at 650 °C in order to induce greater uniformity in the subsequent martensitic transformation. As dem-

Jan 1, 1956

By V. Kondic, W. M. Rumball

HUNT and Jackson' have recently suggested that binary eutectics might be classified according to the entropy of fusion of the component phases. Such a classification was in fact proposed2 some time ago by the authors and the purpose of this communication is to present some supporting evidence. As pointed out by Hunt and Jackson' typical metals of low entropy of fusion, ?S, tend to form normal eutectics (using Scheil's terminology3). Where one phase is metallic (low ?S) and the other nonmetallic (high AS) then anomalous structures are formed. These anomalous structures may either be regular (e.g., Bi-Sn)or irregular (e.g., A1-Si). Finally Hunt and Jackson suggest that where both phases are of high ?S regular faceted structures are formed. In the proposed classification2 the ratio of the entropies of fusion is taken as the quantitative parameter, i.e., ?Sa/?SB = E, where e > 1.0. An extensive literature survey has been carried out2,4 for reliable photomicrographs of eutectic structures. The structures, numbering 151, have been classified into two classes, viz., normal (i.e., lamellar and rodlike) and anomalous. The latter group has been arbitrarily subdivided into regular and irregular anomalous eutectics. Details and sources of the structures are given elsewhere.2,4 In the absence of better thermodynamic data the value of E has been determined for each of the systems making the following assumptions and approximations: l) The entropy of fusion of solid solutions may be calculated additively;5 ii) Inter metallic compounds are treated similarly but with the addition of the following term for fully ordered structures.5 -4.573(Ni lOg10 Nl + N2 log10 N2) where Nl and N2 are the atomic fractions. It might be argued that the entropies of fusion are meaningless as the eutectic phases form at some temperature (T,°K) below their true melting point (Ta or Tb). As AS, = Lo/To if La the latent heat of fusion of a, is assumed independent of temperature, the actual entropy change AS& is given by La/Te. The ratio of entropies of fusion then becomes Fig. 1 shows E' plotted against cumulative totals for normal and anomalous systems from which it is seen that normal eutectics predominate at low values of E'; e.g., some 90 pct of normal systems lie below E' = 1.5 with 65 pct of anomalous systems above this value. Fig. 1 also shows that the regular anomalous structures occur mostly at low values of c' possibly suggesting growth more akin to normal than anomalous growth. The choice of c as a classification parameter was based on nucleation concepts, it being observed that metals widely separated in Sundquist and Mondolfo's8 nucleation series were anomalous and those close together normal. The basis of these classification views is really outside the intended scope of this communication but it is hoped to explain them fully in the near future.7 Having arrived at this entropy of fusion classification by nucleation considerations it is interesting to note that Hunt and Jackson's ideas are based on growth considerations. Elsewhere one of the authors4 has shown that an even better classification of the metallographic structures reviewed is obtained from purely conventional heterogeneous nucleation concepts. It has been found that systems having a value of the parameter +b less than 0.28 are almost certainly normal and above this anomalous: V, is the gram atomic volume. This parameter is nevertheless closely related to c and c', these being an approximation of p. The parameter is also related to the other classification3,8,9 which have met with any success in the past. All these classifications are to be reviewed shortly by the authorsU7 A significant thing about the classifications based on

Jan 1, 1968

By B. A. Kulp, R. R. Reeber

Lattice parameters for the wurtzite form of' CdS mere measured by powder X-ray diffraction techniques over the temperature range 26° to 1000 K'. A negative thermal -expansion coefficient was determined parallel and perpendicular to the c axis below 13.5°K. Near 808°K, the curves of the a, parameter, the differential thermal expansion, and the resistivitv vs absolute temperature exhibited a reversible discontinuity. j\. direct determination of the lattice parameters of CdS over an extended temperature range is of particular use in interpretation of the effects of high pressure on the band structure of semiconductor materials. For example, Langer' has studied the effects of pressure on the absorption edge of CdS from 77°K to room temperature. In order to properly interpret his results it is necessary to know the change in lattice constants as a function of temperature for the wurtzite, zincblende, and rocksalt2 modifications of this material. Thermal-expansion measurements for CdS, wurtzite phase, have previously been determined over a limited temperature range.s74'5 The zincblende phases, of other compounds with similar bonding, have been investigated by Novikova et a! .'' Previous thermal-expansion measurements for metals with hexagonal symmetry have for the most part used bulk sample techniques for low-temperature result. Most X-ray measurements at low temperature have been on materials with cubic symmetry.'jQ The use of X-ray powder-diffraction techniques offers a convenient and accurate method of measurement over a wide temperature range. The X-ray method, besides measuring unit cell size, also provides information about change in structure type and the nature of some of the stacking defects present. I) EXPERIMENTAL PROCEDURE The thermal expansion was measured by X-ray powder-diffraction techniques. Lattice parameters for the range 300" to 1000°K were determined in a 19-cm Unicam camera. From 26" to 300°K a 34-cm back-reflection camera with the Van Arkel" film arrangement was used. The sample was bonded to a copper cold finger below a helium, hydrogen, nitrogen, oxygen, dry ice, or ice brine bath in a metal cryistat. che inner bath was surrounded by another one for radiation shielding. Nitrogen was used in this outer bath for the very low-temperature runs while dry ice and acetone was used for the 195°K point. CuKa X-rays, filtered with a nickel foil, entered the cryostat through a 0.001-in.-thick aluminum window. During a typical 9-hr exposure the sample was continuously rotated through a 30-deg arc. Heat losses from the sample holder limited the minimum temperature attained to 26°K with helium in the inner dewar. Temperature calibration between 25" and 125°K was determined with an Allen Bradley 300-ohm carbon resistor. The resistor was attached to the back of the copper finger with a thin layer of silicone grease. Liquid He, HZ, N2, and O2 baths were used to calibrate the resistor. Temperature variation is estimated to be *2"K with a slightly higher uncertainty at room temperature and 264°K. The thermal expansion of high-purity germanium from 373" to 900°K was measured as a calibration for the Unicam camera. This value was compared with the known value1' for the thermal expansion. A discontinuity in the thermal-expansion coefficient of the CdS a, parameter was observed. Similar discontinuities in the resistivity and differential thermal analysis at the same absolute temperature indicate that the temperature error is *5°K in the temperature range above 500°K. Temperature errors in the Unicam camera below 500°K are appreciably higher due to large gradients in the resistance furnace. II) SAMPLE PREPARATION The CdS was Eagle Picher ultrapure grade. The crystallite size range of the as-received powder

Jan 1, 1965

By S. R. Maloof, W. R. Mitchell, J. A. Mullendore

Preliminary experimental evidence is presented to show that the metallic impurities aluminum, iron, and silicon, and beryllium oxide as found in commercially pure hot-pressed beryllium powder can be reduced to lower Concentrations by Zone -Purification techniques. The reduction in the concentration of aluminum to extremely low levels (10 PPm) is noteworthy, since earlier work demonstrated that aluminum is a major factor contributing to the hot-tearing of beryllium during fusion welding. On the basis of present findings, a method is suggested for producing beryllium metal of improved weldability. EXPERIMENTAL research now being conducted on the purification of beryllium indicates that several of the metallic impurities and beryllium oxide can be reduced to lower concentrations by employing standard horizontal zone-refining techniques. For the preliminary work, commercially-pure hot-pressed beryllium having a purity of 97.84 pct was used. The chemical composition is given in Table I. The hot-pressed beryllium was machined into bars 5/8 by 5/8 by 4 to 6 in. in length. Beryllium-oxide boats which had been conditioned by prior heating in a vacuum of 10"5 mm of Hg at 1100°C for 3 hr were used to hold the beryllium while it was zone melted. The boat with its charge, along with a smaller beryllium-oxide boat containing scrap beryllium to be used for gettering purposes, was loaded into a quartz tube and evacuated to 10-6 mm of Hg. Alter alternately evacuating and flushing three times with purified argon, the furnace tube was backfilled to a pressure of 1/6 atm of A and sealed off by means of a vacuum-tight stopcock. It was then removed from the pumping equipment and mounted on the traveling carriage of the zone-refining apparatus. Heating was accomplished with a 10 kw 450 kc induc- tion heating unit using an induction coil consisting of two turns of 3/16-in. diam copper tubing. Before starting to zone melt, the atmosphere in the furnace tube was gettered by heating the scrap beryllium in the smaller boat to above its melting point and holding at temperature for several minutes. The carriage was then shifted so as to locate one end of the beryllium ingot in the large boat within the induction coil. A molten zone approximately 3/4 to 1 in. was established by loading directly into the ingot. The molten zone was moved very slowly along the length of the ingot. Before each pass was made the procedure for evacuating, flushing, and backfilling, and so forth was repeatedI' Speeds of 1 1/2 and 3 in- per hr were used with up to twelve zone passes. The appearance of a typical zone-refined ingot is shown in Fig. 1. After zone refining, the grains in the ingot appear to be elongated in the direction of zone travel, whilst before, they are essentially equi-axed. A microexamination revealed the presence of a network in the beryllium grains as seen in Fig. 2. Laue spots in back reflection pictures taken of the sample were not sharp, indicating the presence of subgrains in the structure. As can be seen, however, many of the network boundaries cross beryllium grain boundaries. It has been shown by Martin and Moore1 that beryllium solidifies as a bcc structure which then transforms to the hexagonal structure at about 1250oC.It is believed that the observed network is merely the outline of the bcc structure which formed from the melt and are not subgrain boundaries. In order to eliminate surface contaminants, approximately 1/16 of an inch of material was removed from the outer skin of each ingot after zone refining.

Jan 1, 1962

By R. S. Rickard, M. C. Fuerstenau

Anodic polarization curves for iron dissolution and cathodic polarization curves for copper deposition and ferric and hydrogen ion reduction were studied. These results were used to predict the relative rates of the reactions: The rates of the first two reactions were found to be first order with respect to the reacting ion, and the rale-controlling step was diffusion of the ions to the callzode surface. The rate of hydrogen ion reduction in an oxygen-free solution is dependent on the reciprocal of the square root of the hydrogen ion activity. In solutions containing oxygen, the rate of hydrogen reduction becomes dependent on the square root of the partial pressure of oxygen. At equal concentrations copper precipitation occurs about twice as fast as ferric ion reduction. In an air-saturated solution containing 0.01 M cupric ion at pH 2, the rate of hydrogen ion reduction zs about 10 times slower than that of copper precipitation. At pH values below 3, the rate of hydrogen reduction in oxygen-free solutions is very slozu. COPPER cementation has been utilized industrially for many years as a means of recovering copper from aqueous solutions. There are many different types of cementation plants, but all simply provide a means of contacting iron with relatively dilute copper solutions. In the past little or no attention has been paid to rates or mechanisms of the reactions involved. As a result, this and other studies'-3 have been undertaken recently to establish the important parameters in this system. The reactions involved in copper cementation by iron are: Since transfer of charge is involved in every case, these reactions should be considered from an electrochemical point of view. In any electrochemical system involving electrolytic conduction, passage of electrons necessitates ion transfer to the respective electrodes according to Faraday's law. The current in the system will be proportional to the rate of ion transfer, and, therefore, when the current in the galvanic cell is known, the rates of the reactions involved can be determined. Galvanostatic electrode measurements5" are applicable for these electrochemical studies, in that with this technique polarization curves (current density vs electrode potential) are established by impressing a constant current through an electrode and measuring the electrode potential. Once the current density, electrode potential, and reacting area are known for each reaction in this cementation system, the rates of the iron-consuming reactions can be determined. EXPERIMENTAL MATERIALS Several metal electrodes were used in this investigation; they are shown in Table I. Solutions were made with reagent-grade salts and conductivity water. Air, nitrogen, and oxygen were employed. All of these gases were saturated with water and degreased prior to use. EXPERIMENTAL PROCEDURE The experimental work was divided into two general areas. The first area involved galvanostatic measurements, which were used to determine the rates of reactions as a function of various parameters. The second area comprised actual cementation experiments to check the rates predicted from the electrode measurements. Galvanostatic Electrode Experiments. Polarization curves were obtained using a standard galvanostatic technique5'= with the apparatus shown in Fig. 1. The various electrode polarization experiments were run the same way except for changes in solution composition, polarity, and type of measured electrode. The general procedure involved the following steps: 1) a solution was made 0.5 M in sulfate ion and to the desired cation concentrations; this solution was placed in a constant-temperature bath and saturated with the desired gas; 2) the electrodes were cleaned and placed in the electrode holder; 3) the previously cleaned apparatus was assembled and immersed in a constant-temperature water bath (temperature controlled to ± 0.05"C); 4) the electrode was then positioned; the solutions were introduced, and the polarization meas-

Jan 1, 1969

By P. G. Winchell, J. K. Jackson

The strtcture of body-centered Fe-Ni-C alloys (0 to 16.5 wt pct Ni) containing less than 0.015 wt pct C was investigated by measuring the carbon-diffusion peak at low frequencies with a torsion pendulurrr. By determining the Snoek damping in ferrite and in quenched alloys of identical composition, the effects of nickel content and heat treatment were separated, and the existence of a structure in which Snoek damping is greatly reduced was revealed in the quenched ternary alloys. The relative amount of this structure increases with increased nickel content and cooling rate. It is proposed that this structure may be highly imperfect or tetragonal. In Fe-Ni-C alloys containing only ferrite the addition of nickel (0 to 5.1 wt pet) decreases tke area of the carbon-diffusion peak and shifts its position to lower temperatures but broadens the peak only slightljl. The decrease in peak area with increasing nickel implies that short-range interactions with the nickel atoms prevent some of the carbon atoms from participating in the relaxation process. The shift in the peak position indicates a long-range effect of nickel on the movement of carbon atoms so that at 300°K the diffisivity of carbon in -ferrite is increased by about 35 pct by the addition of 5.1 tut pct Ni. The results are presented as distributions of relaxation strength 072 the activation energy under the assumption that the frequency factor is invariant. BODY-CENTERED iron-base solid solutions dissolve carbon interstitially in the octahedral sites. The occupation of an octahedral site produces a tetragonal distortion of the surrounding lattice. These sites fall into three sets according to the orientation of the tegragonal distortion associated with each and may be identified by their lattice coordinates as: type 1/2 0 0, type 0 1/2 0, type 0 0 1/2. At low carbon levels in cubic ferrite the carbon atoms are distributed at random1 among these sites. In high-carbon body-centered iron-base solid solutions, which are unstable with respect to several two-phase mixtures, a tetragonal structure occurs as the product of the diffusionless transformation of austenite. This transformation is commonly accompanied by appreciable microstrain and its product is called martensite. The tetragonality of martensite, a supersaturated solution of carbon in a iron, is maintained prior to decomposition even though local diffusion of carbon occurs as evidenced by carbide precipitation. The tetragonality is taken as evidence that one set of octahedral sites is preferentially occupied by carbon. (This preferential occupation is called ordering below.) Support for this view is presented by the absence of strain-induced carbon-atom diffusion in high-carbon solid solutions (martensite). Such diffusion occurs in ferrite and produces the 40°C internal-friction peak observed at 1 cps frequencies.' Thus, there appears to be qualitative evidence favoring the view that at sufficiently high carbon levels tetragonality arises spontaneously as a result of the ordering of the carbon atoms; however, no experimental evidence exists concerning an equilibrium between cubic and tetragonal body-centered solid solutions. Moreover, attempts to estimate the energy of ordering7-' predict a value too low to maintain the tetragonality observed at medium carbon levels, 0.3 to 0.6 wt pct C. In the present work an attempt using internal-friction measurements is made to establish the presence or absence of tetragonality in low-carbon body-centered solid solutions formed by quenching austenite. Fe-Ni alloys containing about 0.01 wt pct and 0 to 15 wt pct Ni were selected for the study. While the work was in progress Gilbert and Owen10 reported that the transformation which occurs in these alloys during quenching at rates up to 5000°C per sec is not martensitic but is more properly categorized as massive. Thus, the solid solution

Jan 1, 1964

By R. I. Jaffee, C. M. Craighead, G. A. Lenning

THERE has been considerable discussion among A metallurgists and others interested in the development of titanium alloys as to the nature of the fine line markings which appear in the microstruc-tures of unalloyed titanium and a titanium alloys. These markings have been variously called x phase, etching pits, and twins. Figs. 1 to 3 show typical examples of the line markings as they appear, respectively, in high purity iodide titanium, commercial purity titanium rod, and a high purity a alloy containing 0.2 pct N. The commercial purity titanium sample, shown in Fig. 2, contains both line markings and some spheroidal retained-ß phase. The ß phase seen in this micrograph should not be confused with the line marking, since it originates from the iron present as an impurity in the sponge metal. It will be noted that the line markings seem to have a definite orientation pattern with respect to the grains and in titanium noticeably resemble the precipitation occurring in artificially aged high purity Al-Cu alloys. This similarity led to the hypothesis that the line markings observed in titanium and a titanium alloys resulted from precipitation of some unknown phase. With the work of Hägg1 and McQuillan2 as background, there was little reason to suspect hydrogen of being responsible for the line markings. Hagg reported that up to 33 atomic pct H was soluble in a titanium, and McQuillan reported an a solubility of up to 8 atomic pct. Furthermore, McQuillan examined metallographically a 5.1 atomic pct H alloy quenched from 300°C and stated that the structure appeared to be a single-phase a solid solution. Line markings have been observed frequently and were first noted by Jaffee and Campbell3 in 1949. They did not observe them in vacuum-annealed titanium and a Ti-O and Ti-N alloys, but did observe them in Ti-H alloys containing 0.25, 0.5, and 1 atomic pct H. They stated that it was doubtful that the markings were a second phase based on hydrogen, because 1 atomic pct was far from the accepted solid solubility of hydrogen in titanium. Finlay, Resketo, and Vordahl,4 in 1950, concluded that the markings were a true second phase as a result of observations on their etching characteristics. Anderson5 Commenting on these markings in titanium was inclined to associate them with hydrogen, but was loath to call them titanium hydride, because of the reported high solubility of hydrogen. Work at Battelle Memorial Institute for the Navy Bureau of Aeronautics," done concurrently with the work described in this paper, provided information on the experiment converse to that described by Jaffee and Campbell in 1949. This work showed that the amount of line markings in unalloyed titanium could be decreased by vacuum annealing at elevated temperature. It was suggested that this was the result of removal of hydrogen; previously, the data of Jaffee and Campbell9 howed that the markings appeared when hydrogen was added. Recent work at Battelle for the Ordnance Corps, Department of the Army, which is described here, now establishes the identity of the markings as a titanium hydride phase. A sample of high-purity titanium rod 1/4 in. diam, 0.009±0.001 wt pct H (0.43 atomic pct), Fig. 1, was resistance heated in a 0.05 micron vacuum for 41/2 hr at 1500°C. During the early states of heating, the pressure increased to 150 microns, indicating that gases were being evolved. After cooling in vacuum, the iodide titanium rod was analyzed for hydrogen. Vacuum-fusion analysis showed 0.0006 ±0.0003 wt pct H (0.03 atomic pct) to be present.* Metallographic examination of the vacuum-annealed rod showed only a transformed ß structure. There was no evidence of fine line markings, Fig. 4. A section of this vacuum-annealed rod was cold rolled to 0.040 in. thick sheet and annealed for 1 hr

Jan 1, 1953

By R. A. Graham, R. W. Rohde, J. R. Holland

In a shock wave compression study of a martensitic Fe-30 pct Ni alloy, Graham, Anderson, and Holland' found a region of unusual compressibility extending from a few kilobars up to about 20 kbars. This region was characterized by a continuously decreasing shock wave velocity with increasing stress. In this note, we report the results of a study initiated to determine the source of this unusual compressibility. In earlier shock-loading experiments on martensitic Fe-30 pct Ni, Loree et a1.' determined that a phase change occurred at about 80 kbars. Subsequently, Bowden and elly found that samples of a similar material shock-loaded at about 100 kbars experienced a martensite-to-austenite transition. Their crystal-lographic analysis showed the transition was a direct reversal of the established martensite system to the austenite phase. The shock loading in the present experiments was accomplished by precisely controlled plane impact between two discs of the sample material. This configuration created a state of uniaxial strain in the two the front of a projectile which was accelerated to various velocities. Planarity between the impacting discs was controlled to 5 x 10-4 rad, and impact velocities were measured to 0.5 pct. From a knowledge of the Hugoniot relationships of the material and the impact velocity, the shock-wave stress can be calculated. Samples were shocked at 18, 50, 70, and 100 kbars. The specimen discs shocked at 18 and 50 kbars were instrumented with quartz gages to observe the stress time profile. A more detailed description of the experiment has been given previously.4 The shock-loaded specimens were recovered intact by decelerating them in a catcher filled with soft wood. Discs of Fe-30 pct Ni nominally 12.7 mm thick and 64 mm in diam were cut from bar stock. Chemical analysis of the material provided the following composition in weight percents: 29.5 Ni, 0.50 Mn, 0.21 Si, and 0.10 C, with traces of Cr, Cu, Mo, Al, and Co. The discs were austenitized at 650° C for 2 hr at a pressure of less than l0-' torr, and furnace-cooled. Martensite was then produced by quenching the discs at -196" and holding 168 hr. The initial martensitic structure formed is shown in Fig. l(a). Typical microstructures of samples shock-loaded at 18 and 100 kbars are shown in Figs. l(b) and l(c). The banded structure with the austenite regions running parallel to the shock propagation direction was characteristic of samples shock-loaded at all stresses. No such banding was found in samples that were not shock-loaded. Comparison of Figs. l(b) and l(c) shows that, although under the optical microscope the two microstructures are quite similar, the martensite plates of the specimen shocked at 100 kbars etched in a distinctly different manner. This unique etching behavior of martensite in samples shock-loaded above the 80 kbar transition point has also been observed by Leslie el al.' These investigators determined from X-ray and electron microscope studies that fcc martensite, similar in appearance to normal bcc martensite, had been formed under shock loading. A quantitative determination of the amount of austenite present after shock loading was made by X-ray techniques. Both the methods of Lindren' and Miller' were utilized. The results are shown in Table I. The error limits are those experimentally observed from two independent analyses of each sample and represent the precision of the experiment. The absolute error, which should be similar for each sample, is believed to be about 10 pct. It is seen that as the shock stress increases the amount of austenite present also increases. On the basis of the X-ray and microscopic observations, we conclude that shock loading below the 80 kbar ersal will initiate a martensitic reversal: the amount of reversal that occurs increases with in-

Jan 1, 1969

By H. A. Fisch, A. J. Opinsky, J. L. Bartos

Fort?-six- two-grain bounduries were studied in doped tungsten 1amp wire that had been heated to 3450°K very vapidly. Back-reflection Laue photograms were taken of the grain boundary partners and solved, using a computer. The resulting distribution of large-grain orientations was analyzed. Rotational axes were calculatedfor each grain boundary pair. The axes of least rotation show an anisotropy of rolalional angle. WhILE the orientations of large grains in straight doped tungsten wire have been studied extensively, little has been done to characterize the boundary that separates them. Rosi.1 upon seeing the jagged interlocking boundary in doped tungsten wire. inferred that the grains should have a preferred orientation, and proved it. He interprets his data as indicative of a single ideal orientation—(531)— being parallel to the wire axis. Rieck2, 3 and Mannerkoski' have essentially confirmed this finding. Opinsky, Orehotsky, and Siegle argue that the data suggest that (211)—a plane-is parallel to the wire axis instead. This (211 ) contains the (531). and hence would appear to be a more general case.5 Ahlborn and wasserman6 report that their orientation data stretch from (110) through (531) to (311). This path is the locus of {211 Recently. Opinsky has suggested that the data can be interpreted as a series of ideal orientations that can be related by coincidence relationships to the (110) fiber texture. All these investigators have determined the orientations of the large grains by the back-reflection Laue technique, and they have all obtained similar results. Only Rieck5 as methodically looked at the large grains along a wire. He was apparently not concerned with the relationships of the large grains across their common boundary, but only with their individual orientations. This research was undertaken to fill this gap in the under standing of the mechanical properties of tungsten wire. Only the crystallographic aspects of the problem will be reported here. In addition, the large number of large-grain orientations, which was accumulated in this study. will also be analyzed. EXPERIMENTAL MATERIAL AND TECHNIQUES All of the data were obtained from four pieces of the same 7-mil-diam General Electric tungsten lamp wire. Twelve-inch lengths of the cleaned wire were heated in vacuum by the flick of a switch to 90 pct of the amperes necessary for fusion (a temperature of approximately 3450°K) and held for 1 min. On the basis of other data. it is estimated that the rise time is of the order of 0.5 sec. The four pieces of wire could have come from anywhere within about 18 ft of wire. Consecutive segments of each piece of wire were mounted in bakelite and polished metallographically The polishing procedure followed that described by woods. "nly boundaries between two grains that each occupied the entire cross section of the wire were photographed and X-rayed. The metallographic sample was mounted accurately with respect to the X-ray unit. Steps were taken to read the films as accurately as possible with the conventional Greninger net. The data were then entered into computer programs that solve the Laue film by precisely the same techniques that are used in the solution of the film by means of the stereographic projection. The use of the computer has probably shifted the main source of error from the stereographic plot to the reading of the film. On the basis of several experiments, it is thought that the error associated with the entire process, from mounting the sample through the computer solution, is of the order of about one-half degree. The output of the computer is the direction cosines of the fiducial direction (the wire axis) and the locations of the three cube axes. The former data were used to calculate the differences in orientation between the two grain boundary partners, and the latter were used in the calculation of the rotational axes that could bring the two grains

Jan 1, 1970

By N. Jaffee, J. E. Dorn

An activation energy of 27,400 5 1000 cal per mole was obtained for the creep of poly crystaLline aluminum over the temperature range of 273° to 350°K, at strains varying from 0.003 to 0.230. Stresses varied fronz 2250 to 6000 psi, strain rates from 0.1145 x 10-5 to 29.5 x 10~s min.-1. A strain temperature-compensated time correlatzon was oblained indicating that the structure was depedent solely upon the strain at a given stress; X-ray diffraction analyses and tensile tests on precrept specimens verified this. PREVIOUS investigations1 have revealed that the apparent activation energies for creep of high-purity polycrystalline aluminum are insensitive to the applied stress, strain, and strain rate. On the other hand, the apparent activation energy is dependent on the temperature1,2 as shown by the data reproduced in Fig. 1. Each plateau of this curve is associated with a unique rate-controlling dislocation mechanism for thermally activated creep. It is now very well-established that over the high-temperature range, where the activation energy is about 35,500 cal per mole, the creep rate is controlled by the dislocation climb mechanism. 3 Over the intermediate range of temperatures, from about 260° to 370°K, another unique activation energy, namely 28,000 cal per mole, is obtained. Current evidence strongly supports the contention that cross slip is the rate controlling mechanism for creep in this case: 1) Crystal recovery can occur by a number of mechanisms prominent among which are climb and cross slip. Astrom,4 using the Borelius micro-calorimeter, reported rapid crystal recovery of cold-worked polycrystalline A1 at 340° to 370°K giving an activation energy of 28,000 cal per mole and additional recovery at 455° to 478°K giving an activation energy of 3 6,000 cal per mole. Whereas the latter recovery process was clearly associated with the dislocation climb mechanism, the former, 28,000 cal per mole process could have arisen as a result of the cross-slip mechanism. 2) Single crystals of high-purity A1 favorably oriented for octahedral glide exhibit an activation energy for creep of about 3400 cal per mole over the range from 0° to 400°K and that of about 28,O:O cal per mole over the range from 600° to 775°K . Whereas the lower activation energy process is accompanied by the development of sharp slip-band traces, the 28,000 cal per mole process was uniquely characterized by extensive cross slip. 3) Theoretical calculations by Seeger and schoeck6 suggest an apparent activation energy for cross slip in A1 of about 24,000 cal per mole and another theoretical estimate by Friedel 7 gives the observed 28,000 cal per mole as the theoretical activation energy. It is therefore consistent with this evidence to, at least tentatively, ascribe the 28,000 cal per mole creep process in A1 to the operation of cross slip as the rate controlling mechanism. Although creep has been rather well-explored in the dislocation climb range, very little, excepting perhaps the activation energy, is known about creep in the cross slip range. This investigation was undertaken, therefore, to reconfirm the activation energy for creep in polycrystalline A1 in the intermediate temperature range and to study the strain-time relationship for creep at a constant stress in the cross slip region. EXPERIMENTAL PROCEDURE AND TECHNIQUE The three types of high-purity A1 alloy specimens described in Table I were used in this investigation. Specimens were milled from 0.100 in. thick cold-rolled sheets with their tensile axes in the rolling direction following which they were given a recrys-tallization and grain growth heat treatment at temperatures well above those which were subsequently employed in the creep tests.

Jan 1, 1962

By F. D. Patchen

Changes in the properties of partland cement upon the addition of fine-ground silica are discussed. Data were collected from formulations cured for periods up to 60 days at temperatures varying from 180" to 350°F. Compressive strengths, permeabilities and thickening times are compared. It is demonstrated that addition of find-ground silica in proper proportion results in a hydrated cement with superior properties above 230°F; i.e., early compressive strengths are higher and increase with additional curing, permeabilities are significantly lnrver and thickening times are increased. X-ray analysis of hydrated cement samples have been used to determine more fully the complex chemistry of these systems. Dicalcium silicate alpha-hydrate was found in all neat cements cured above 260°F. A 10 weight per cent silica addition produced a maximum amount of this phase. Xonotlite and tobermorite have been identified as hydration products at 320°F in systems containing 35 or more weight per cent silica. Formation of these components appears to be primarily responsible for the gain in strength and reduction of permeability observed after the addition of silica. INTRODUCTION For several years it has been recognized that loss of compressive strength occurs when portland cement is cured at a temperature above 250°F. This phenomenon has been called "strength retrogression". More recently it has been recognized that increased permeability usually is associated with decreased compressive strength at these temperatures. At curing temperatures above 300°F, these factors reach serious proportions. A compressive strength loss of 50 per cent between curing times of 24 hours and 14 days has been reported,' and permeabilities as high as 10 md have been measured in neat cements cured at 320°F for seven days.' Earlier investigators,' although unsuccessful in eliminating strength retrogression, stated ". . . .it appears possible that by the proper selection of additives for cement, strength loss with age at high temperatures can be either eliminated or considerably lessened". One solution suggested was the use of a pozzolanic material, hydrated lime, and a retarder. Menze15 in 1934 discovered that portland cement containing finely divided silica possessed high compressive strengths when the curing temperature was 350°F. The use of such a composition is now an established practice in the manufacture of steam-cured cement blocks. It is reasonable to assume that the same beneficial effect can be used advantageously in high-temperature oilwell cementing operations. Ludwig first mentioned the use of sand for this purpose. Later, a portland cement containing moderate additions of silica flour and retarded with CMHEC, a cellulose derivative. was recommended for high-temperature application." These formulations have been successfully employed in the field."'" This paper presents the results of a research program to test the efficacy of finely ground silica in increasing the compressive strength and reducing the permeability of oilwell cement compositions containing portland cement. Included also are X-ray diffraction patterns of powdered samples which have been used to determine some of the important chemical reactions occurring in these systems. LABORATORY PROCEDURES AND MATERIALS All compressive-strength test specimens were prepared and cured from a slurry in accordance with API RP 10B, with two modifications. A final curing pressure of 5,000 psi was maintained for a one- or three-day period; but, for further curing, pressures were reduced to the vapor pressure of water at the curing temperature. Also, the water content was varied in samples containing silica to maintain the slurry viscosity within a range of 5 to 15 poise as measured on a consistometer. Test specimens at 180°T were cured in a water bath at atmospheric pressure. At higher temperatures the specimens were cured in a steam-jacketed high-pressure curing vessel at 5,000 psi for a one- or three-day period. cubes cured for longer periods were removed from the pressure-curing vessel after 24 hours, sealed in small bombs containing water and placed in an aging oven.

By C. R. Simcoe, T. J. Koppenaal

The effect of direct and alternating currents on the quench aging behavior in an Al-4 pct Cu alloy has been investigated by use of resistivity measurements. The reaction rate increases with direct current density in the investigated range of 1000 to 3000 amps per sq cm. Alternating current of a specific frequency is observed to decrease the reaction rate by a factor of about two. ErDMANN-JESNITZER et al.,1,2 investigated the effect of direct and alternating current on the quench and strain aging behavior of low-carbon iron. In both cases they observed that direct current accelerates the aging reaction and that an alternating current can completely inhibit the aging process. The purpose of the present research was to investigate these phenomena in the quench aging behavior of an Al-4 pct Cu alloy and to determine the correlation, if any, between the increase in the observed reaction rate and possible heating effects accompanying the aging with direct current. EXPERIMENTAL PROCEDURE The aluminum-base alloy used in this investigation contained 3.98 wt pct Cu (1.72 at. pct) with impurities of 0.003 pct Si, 0.002 pct Fe, and 0.001 pct Mg. The alloy was rolled into a ribbon about 0.36 mm thick and 3.7 mm wide. The ribbon was split end from end with the ends being used as electrical contacts and the central unsplit section, about 5.3 cm in length, serving as the specimen. Solution treatment was accomplished by quenching from a salt bath at 520° ± 2°C into water at 25° ±0.5°C; the specimen was immediately transferred to a liquid nitrogen bath. All aging was done at 75° ± 0.2°C in a rapidly stirred ethylene glycol bath. Resistivity measurements were made by the double potentio-metric method employing a standard resistor; for purposes of convenience, the resistance measurements were made at 77°K. The transfer of the specimen from the aging bath to a liquid nitrogen dewar was accomplished in less than 1 sec. The aging kinteics were studied with the specimen being aged under the influence of direct current, alternating current, and without any type of current. The amount of heating caused by direct currents was investigated by soldering a thermocouple on the specimen surface and measuring the temperature change after applying various current densities. The temperature was observed to increase (from 75°C, the aging temperature) about 0.8°, 3.0°, and 8.5°C with direct currents of 1000, 2000, and 3000 amp per sq cm, respectively. When employing alternating current, the circuit was monitored with an oscilloscope and adjustments made to produce a normal sine-wave pattern. RESULTS Prior investigations3'4 in the aging characteristics of A1-4 wt pct Cu have shown that the resistivity initially increases, reaches a maximum, and then decreases to values significantly lower than the as-quenched resistivity. This type of behavior is seen in Fig. 1 which shows the resistivity at 77°K as a function of aging time at 75°C observed in this investigation. At this aging temperature, the resistivity maximum is seen to occur in about 1 to 1-1/2 min and it further decreases to the as-quenched value in about 24 min. The behavior observed with the application of direct currents of 1000, 2000, and 3000 amp per sq cm during aging is shown in Fig. 2. With increasing current densities, (+) ?pmax decreases and the rate of the aging process increases. The time necessary to reach the as-quenched resistivity value decreases from 24 min for aging without current to about 16, 10, and 7 min for aging with 1000, 2000, and 3000 amp per sq cm direct current, respectively. Fig. 3 shows the results obtained when two changes were made during the aging with 2000 amp per sq cm

Jan 1, 1963

T.P. 1958 D. H. DAVIS.*-Mr. Carris has ably pre¬sented the problems of coal preparation, particularly those brought about as the result of increased mechanical loading. Many mines are now crushing the run-of¬-mine raw coal through 12 or 15 in. immediately after it is dumped from the mine car and before it is handled by conveyors. At certain mines the removal of large slate mechanically has been solved by the installation of scalping screens consisting of bars or rails set at spacing of 12 to 15 in. on the top deck of the main raw-coal sizing screen. At the plants where these are in use, the plus 12 or I5-in. raw material contains 60 per cent or more free rock or slate, so that it is easier to pick coal than to pick rock. After the screening at 12 or 15 in. is accomplished, the oversize material is diverted to one side of the shaker, where a picker allows the refuse to pass off the shaker into a chute and guides the coal back onto the shaker. After the large rock has been removed, it is then advisable to reduce the size of rock by crushing before putting it onto a conveyor. A great saving in picking labor has been made by the use of such scalping screens, which in addition has maintained the plant capacity when handling a dirtier feed. The great increase in percentage of coal under 1/4 or 3/8 in., coupled with a greatly increased percentage of free impurities in this size, is the most serious problem that has been brought about by mechaniced loading. Many plants that formerly prepared good quality 1/4-in. to o or 3/8-in. to 0 slack by air¬cleaning are faced with installing facilities for wet-cleaning mechanical and heat-drying and water clarification. The change to wet prepa¬ration of fines usually affects the marketing problems for many coal companies. There will also be problems involving wastage of fine solids and stream pollution. Much work will be necessary for improvements in present practice. Pneumatic and electrostatic methods for cleaning fine coal should receive consider¬able attention in an attempt to avoid wet¬cleaning of fines. The preparation supervisor and the plant personnel should not be satisfied merely to run the cleaning-plant machinery and assume that the job will be done for them. Coal is a heterogeneous substance and in many instances is constantly changing. This is more true when full-seam methods are in use. Only persons trained to observe closely and make the necessary changes in plant operation will be able to solve the many problems that will be imposed on the preparation plants. The plants will require skillful and observant operating men. With full-seam mining methods, seams of roof coal will often be mined. Under the former mining systems these strata were not mined and frequently they contain a high per¬centage of bony and banded partings. The mining of this coal will greatly alter the washability and specific gravity characteristics of the feed to the plant. Often the percentage of 1.40 to 1.60 sp. gr. increased from 3.5 to 7.5 per cent, or nearly a 1.0 per cent increase in the ash content of the 1.60 sp. gr. float coal. Thus the cleaning problem is much more difficult, and it is unlikely the same ash content of cleaned coal can be produced as before without a decrease in recovery or the produc¬tion of a "middlings" product. Mr. Carris had this in mind when he said "The quality of the

Jan 1, 1946

By T. Szaplonczay, D. E. Light

The auto-oxidation of sulfur dioxide to form sulfuric acid in commercially produced barren and recycle liquors, containing 1.66 and 4.56 gpl Fe respectively, from Elliot Lake was studied. A gas mixture, containing 3.5% SO2 and air was introduced into these liquors through a frit at a rate of 6.0g SO2 per hour per liter of solution. The barren and recycle were heated to 60°C. At the end of the first 24 hours of the process the concentrations of acid produced were 158 and 220 gpl H2SO4 for the barren and recycle respectively. Corresponding conversion efficiencies were 70 and 83%. In the period from 24 to 36 hours the rate of acid production decreased so that the H2SO4 concentrations at the end of 36 hours were 180 gpl for the barren and 260 gpl for the recycle; average conversions for the initial 36 hours were 55 and 68% respectively. Many investigators1'2'3 have considered the oxidation of sulfur dioxide to form sulfuric acid in aqueous solutions containing metallic ions. A summary of the work performed on this subject prior to 1946 has been presented by Keyes4 who described the auto-oxidation process as follows. Keyes also suggested that in this second reaction ferric ion serves as a catalyst for the conversion of sulfur dioxide to sulfuric acid. Neytzell-de Wilde and Taverner5 have reported on the use of auto-oxidation to produce a 4 to 5% sulfuric acid solution for use in a cyclic leach process at South African uranium plants. This paper describes investigations carried out to determine the feasibility of using the auto-oxidation process to produce sulfuric acid solutions of as high concentration as possible using commercially produced uranium plant barren and recycle liquors from the Elliot Lake region of Ontario. These liquors contained 1.66 and 4.56 gpl respectively of iron, chiefly in the ferrous state. This investigation is of particular interest to the uranium producers in Elliot Lake where the ores contain 3 to 4% sulfur, as sulfides. Flotation of these sulfides and roasting of same to produce SO, bearing gases have been successfully demonstrated using conventional methods.6'7 Barren and recycle solutions, as described previously, are also produced in large tonnages by these uranium producers. The experimental method and apparatus used to examine the auto-oxidation process will be described first. Then the tests carried out on laboratory prepared aqueous solutions containing ferric sulfate to determine optimum conditions will be presented. The next section will describe the tests using the commercial barren and recycle liquors. A final section will summarize the conclusions of the work. EXPERIMENTAL APPARATUS AND PROCEDURES Two types of solutions were used in the tests, a laboratory prepared aqueous solution containing definite amounts of ferric sulfate, and commercially produced iron bearing solution obtained from a uranium leaching plant in the Elliot Lake area. Two liters of a solution to be examined were placed in a 3 in. diam by 30 in. long glass cylinder with a cross-sectional area of 7 sq in. A mixture of air, and SO2 supplied from a lecture bottle for convenience, was introduced into the bottom of the vessel through a standard 30 mm glass frit, 700 sq mm in area. The temperature of the apparatus was controlled by a heating tape and variable transformer. Exhaust gases were vented to atmosphere. At the end of the test period the gas flow was shut off and the solution was analysed for H2SO4, Fe ++, and Fe +++, by conventional means. Four variables were examined using the laboratory prepared aqueous solution of ferric sulfate. First, the effect of SO2 concentration, from 1 to 1076, in the feed gas mixture was considered. Second, the temper-

Jan 1, 1964

By J. A. Roberson, A. M. Adair

A plate of an 18 pct Ni (250 grade) maraging steel was homogenized, cold rolled, and re crystallized to produce a fine grained, homogeneous structure. The strength properties were then measured as a function of cold rollinq prior to aging. Material in the solution heat treated condition exhibited substantial strengthening from cold rolling. The strength increase and the rate of strengthening were proportional to the amount of cold work. This benefit, of up to 50 ksi, was retained during aging for 3 hr at 900°F. The structures produced by cold rolling and aging have been systematically investigated by transmission electron microscopy. The dislocation debris produced by cold work influences the morphology and distribution of precipitate particles as well as the rate of precipitation. An essential feature of the high nickel maraging steels is a martensitic transformation from y to a with the resultant production of a supersaturated bcc solid solution. This solution decomposes to form precipitates during a subsequent aging treatment.' The martensitic transformation in low carbon Fe-Ni alloys and the mechanical properties of these alloys have been studied extensively.2"4 The carbon content of maraging steel is kept very low to avoid the embrittling Fe3C phase and to minimize the acicular, or twinned martensite common to high carbon steels. Maraging steels are usually strengthened only by heat treatment, but cold working is also quite effective in this respect as was shown by Decker, Eash, and Goldman,' Kula and Hickey,5 and the present work. There is, however, some disagreement between the present results and those of Kula and Hickey. Because of their unusual combination of strength and toughness, the maraging steels are suitable for applications where reliability at high stress levels is required. The strength level is achieved through a combination of solid solution hardening,' substructure hardening,3 and precipitation hardening,6 of which the last is the most important. The toughness of these steels has been reported to be a result of the absence of brittle grain boundary precipitates,7 and the fact that the matrix is a massive martensite rather than a twinned martensite.' One of the problems in the manufacture and use of maraging steel has been its tendency to develop laminations in sheets parallel to the rolling plane.9,10 The present investigation was intended to determine whether or not an improvement in the microstructure of this steel by thermomechani-cal processing would improve mechanical properties. I) EXPERIMENTAL PROCEDURES A plate ($ in. thick) of an 18 pct Ni-250 grade maraging steel* was obtained and examined metallo- "Analysis, wt pct: 18.1 Ni, 8.6 Co, 5.3 Mo, 0.4 Ti, 0.07 Mn, 0.05 Cr, 0.05 Si, 0.002B. 0.1 Cu. 0.05 Al. 0.002 Zr, 0.016 C, balance Fe. graphically. It was found to contain laminations as described by previous investigators. This plate was then annealed at 2100°F for 24 hr to promote homogeneity. Pilot studies on small coupons had shown that this was the minimum temperature required to accomplish homogenization in 24 hr. Since this treatment caused excessive grain growth, the steel was subsequently cold rolled and recrystallized. Cold rolling was accomplished at about 5 pct per pass to a total reduction of 66 pct. The standard solution annealing treatment (1 hr at 1500°F) was adequate to cause recrystallization of the cold worked steel. The recrystallized steel was then cold rolled and ground by various amounts to produce specimens with uniform thickness and selected amounts of cold rolling. All tensile specimens with less than 90 pct cold reduction were ground to a thickness of 0.050 in. Specimens with 90 pct cold reduction were 0.025 in. thick. All

Jan 1, 1970

By N. J. Grant, B. C. Giessen, U. Wolff

With the aid of the splat-cooling technique of rapid quenching, metastable alloy phases based on antimony ad bismuth have been prepared. At room temperature, simple cubic phases were found in the Sb + Au, 56 + Yd, and Sb + Ni systems. Another phase was retained in the Au-Bi system at -190°C; in this systent, a microcrystalline phase and a rhombohedrally distorted phase were also observed. The crystal chemistry of these phases is discussed in terms of their valence electron concentration range. A remarkable five-step Phase sequence was found upon heating and equilibration of certain metastable Au-Bi alloys. USING the splat- cooling technique of rapid quenching from the liquid,' a large number of metastable intermediate alloy phases have been prepared in recent years. Thus, Hume-Rothery phases of the hexagonal type based on the valence electron concentration have been produced in systems where they are missing in equilibrium, e.g., Ag-Ge 1 or AU-Sb. 3 In a comprehensive study of metastable alloys based on tin, simple hexagonal metastable y phases were found, raising the total of stable and metastable y phases from 3 to 13. 4 In the In-Bi system, five new metastable intermediate phases were retained.' In most cases, the new metastable phases have some crystal chemical characteristics in common: a) They have rather simple, elementlike structures with few atoms per unit cell, e.g., y phase,4'5 or the metastable primitive cubic phases in the Ag-Te and and Au-Te systems.= b) Their structures are often related to those of the terminal phases, or to elements that are straddled in their position in the periodic table by the constituent elements; this is the case for metastable ß (In-Bi) with the structure of white ß Sn,' or metastable a (Cd-sn). 4 c) They seem to represent the crystal structures appropriate to certain average group number (AGN) ranges in which no or few single-phase structures exist in equilibrium. These average group number ranges can probably be equated in the case of many elements of the Ba(Zn)-Bs(As) groups with the averaged valence electron concentration. In the present work, the splat-cooling technique has been applied to the alloys of antimony and bismuth. In equilibrium, neither element forms intermediate alloy phases with other B elements from groups Bl(Cu) to B4(Si) that contain more than 66.7 at. pct of Sb or Bi, respectively, e.g., AUSb 2. 7 Thus, antimony- or bismuth-rich phases, which might be closely related to antimony or bismuth, and which would reflect struc- tural changes occurring on alloying with small amounts of addition elements, do not exist. Also, the terminal solubility of antimony and bismuth for all B metals is very small. It was therefore of interest to investigate the influence of enforced, nonequilibrium solid solution of addition elements in antimony and bismuth on their crystal structures. It was found that additions of group B 2-B4 metals act distinctly differently than additions of group T 10 and B1 metals. The latter set of additives will be treated here. EXPERIMENTAL TECHNIQUES Alloy preparation and splat-quenching followed closely the procedures described in Refs. 1, 4, and 5; they consisted of the preparation of master alloys from high-purity starting materials (99.99+ pct) and blast-atomizing of 20-mg quantities onto copper sheets. The resulting approximately 1 to 5 u thick foils were investigated by X-ray diffraction. The cooling rates obtained are of the order of 107 to 10 8 C per set,'" causing solidification to proceed at a rat; of approximately 100 cm per sec.' Powder patterns of the resulting fine-crystalline materials were taken with filtered CuK, radiation. For antimony alloys the substrates were held at room temperature and for bismuth alloys at -190°C both during splat foil deposition and X-ray investigation. Annealing treatments at several

Jan 1, 1969

By U. Andres

U. Andres (The Technion, Haifa, Israel)-The authors are to be congratulated on their research, new and interesting, in a field which, however, the present writer and his assistants have been trying since 1962 to bring to the attention of the scientific and technical world, by means of publications covering several different aspects of this many-sided problem. We announced, in these publications, the results of magneto-hydrostatic (MHD) separation of magnetic and nonmagnetic minerals which we made in magnetic fluids (in paramagnetic solutions and suspensions of magnetic particles) which interact with permanent" and impermanent" inhomogeneous magnetic fields. There is no doubt that the authors have succeeded in creating ferrimagnetic colloids made by dispergation of 10-30 wt % of magnetite in kerosene in the presence of 7-10 vl % oleic acid. It may be judged by [Figs. 3 and 4] that the suspensions represent great stability systems and also "hysteresisless" fluids [(Fig. 5)] which bear witness to the existence of superparamagnetic suspensions. If the viscosity of these colloids is not so high, and the cost of preparing them can be covered, the fluids which have here been received can possibly be of interest to the technology of mineral separation. The potential possibility and technological perspective of similar media is demonstrated in our joint work with G. Bunin18 in 1965. The research carried out on the pushing out of cylinders and spheres from magnetic fluid is impressive; it would, however, have been interesting if the authors had compared their results with those on the pushing out of bodies from magnetic fluids adduced from our work." In my books" there happens to be described an experiment on the pushing out of a sphere in magnetic wedge field, in which the pushing out force of suspension with 5 vl % of magnetite reached 125 g per cc. This testifies to interesting possibilities of these media. Our works' shows also, on the basis of Laplace's equation the receiving of an interesting distribution of pushing out forces by selection of corresponding configurations of magnetic poles. Objection must be made to the authors' introduction of the term "antigravity force." The direction of this force does not depend on the direction of the gravity force and their opposition, i.e., antigravity is only one particular case. I am in full agreement with the authors' statement that magnetic pushing out forces "can be utilized to separate nonmagnetic objects according to variations in their densities." It is, however, necessary to make only one important addition, which can be found in our research on MHS at the Technion in Haifa, namely, that magnetic fluids in magnetic fields can be utilized in the separation of minerals according to variations in their magnetic susceptibility. The authors' conclusion is especially important considering the fact that in our publications" (including one in English") some designs of MHS separators and some technological results received on MHS-separators have already been described. S. E. Khalafalla and G. W. Reimers (Authors' Reply)We were aware (mostly through Chemical Abstracts of the American Chemical Society) of the use of solutions of paramagnetic salts such as manganous sulfate in agnetohydrodynamic separations by the Russian school headed by G. Bunin."n Professor Andres' contributions to this school are also evident in his application of the ponderomotive forces in an electrolyte placed in electric and magnetic cross-fields to separate coal and other useful minerals." It must be realized, however, that the foregoing research by Bunin et al." involved the use of magnetic fields and field gradients of very large magnitudes, i.e., such that H • grad H - 101 Oe2 per cm. The magnetization of paramagnetic salt solutions range from 0.1 to 0.5 gauss (at fields of a few thousand Oersteds) and never reach saturation. The magnetization of our superparamagnetic ferrofluids ranges from 100 to 500 gauss and reach saturation at fields of a few hundred Oersteds. Hence, to effect the same separation, the quantity "H

Jan 1, 1974

By Donald L. Johnson

REVIEWS on the applications of optical ellipsometry to the study of polycrystalline metal surfaces and oxidation-corrosion processes have been presented by Kruger and Hayfield, and white.' In view of the ability of this method to measure differences in film thicknesses to below 10A, ellipsometry was used to study nitrogen chromium surface reactions during the early stages of growth. Chromium specimens in the form of a vapor-deposited film on glass substrate were exposed to purified nitrogen for selected time increments at 700°F. This temperature was selected because it gave the best combination of exposure time and thickness change. Ellipsometer readings were taken immediately after exposure to determine polarizer and analyzer angles using monochromatic light at 5000A.2 After each set of readings was obtained, the specimens were exposed for an additional increment of time in nitrogen and a new set of readings taken. Film thickness was computed from LUCy's3 equation for nonabsorbing thin films. In Lucy's derivation, two equations for computing thickness were obtained. Peterson and Bashara4 have discussed the validity for using one of these equations when an independent value of the refractive index is used. The film was assumed to be Cr2N (Ref. 5) and was determined to be optically isotropic. Optical isotropy was determined by rotating the films about the surface normal and noting that the azimuth and phase difference did not change. Since Cr2N is reported to be hexagonal, it is assumed that optical isotropy results from randomly oriented grains of very small size. The thickness was computed using refractive indexes varying from 1.6 to 2.0. The maximum variation in computed thickness over this range for a given exposure was 20 pct. The data reported are based on a refractive index of 1.8. Reference readings on the metal surface were taken just prior to nitrogen exposure. No attempt was made to deoxidize or denitride the metal surface after vapor deposition, although specimens were observed to gain approximately 2A per day in air at room temperature. Twelve experimental exposures were made on a single sample over a total accumulative time of 132 min. Additional experiments over longer time intervals were made to determine the effect of heat up and cool down during and after exposure. As a result of these experiments, the thickness computed for each exposure was increased by 4A. Time profiles based on the linear law are similar to the oxidation profiles for many metals.' The rate of accumulation of the nitride film is rapid at first and slows down as the film thickness increases. The parabolic law x2 =k1t +k2 [11 gives a straight line in Fig. 1; therefore, it is concluded that the kinetic behavior is parabolic. Sey-bolt and Haman5 reported similar results for CrlN on chromium at 1000°C although the order of the thickness was much greater (10-3 cm). Analysis of thickness data based on other values of refractive index shows that the curves are shifted vertically but retain the same shape. According to Mott,7 it is expected that for thin films k2 is zero in Eq. [I]. Since the plot in Fig. 1 is linear through the origin, k2 is zero and the parabolic nature of the film growth is further indicated. The parabolic rate constant, calculated from the slope in Fig. 1, is 4.4 x 10-16 sq cm per sec. The corresponding work (+) required to bring an electron from the metal into the oxide is estimated to be 1.5 ev.7 The author thanks professor N. M. Bashara and D. W. Peterson, Electrical Materials Laboratory, University of Nebraska, for making the ellipsometer available and providing technical assistance. The Electrical Materials Laboratory is supported by the office of Naval Research, NONR 3718 (01). This work was partially supported by the Engineering Experiment Station and the Mechanical Engineering Department, University of Nebraska.

Jan 1, 1967

By William F. List, Marvin A. Schuster

Monolithic electrooptical mosaics of 2500 photo-transistor elements with internal row and surface column interconnections have been fabricated by epitaxial-planar diffsion techniques. Unique access to any elemed of the mosaic is provided through one of the fifty X and one of the fifty Y external leads. The sensor mosaic can be used for imaging applications and has a 1:1 aspect ratio, a resolution of 100 lines per in., a dynamic range in excess of lo3, a sensitivity of 0.5 per pw, element response within a 2:l range over 75 pct of the mosaic, and a minimum detectable signal of the order of 10-' w. Design considerations which ilnpose particclar fabrication yequirements are disctussed along with some of the problems peculiar to the epitaxial and planar-dijfision techniqes by which this device is fabricated. The solid-state silicon monolith has an imaging area of 0.5 by 0.5 in. and an element center -to- center spacing of 0.010 in. J monolithic solid-state electrooptical device has been developed as the photon-sensing mosaic for use in an image-converter system.' This system is the solid-state equivalent of the vidicon and can image radiation in the visible and near-infrared regions of the spectrum. The electrooptical image converter has the desirable size, weight, reliability, and low-power characteristics inherent in solid-state devices. It is read out by direct-wire means rather than by beam-scan techniques. The packaged unit is shown in Fig. 1. A selection of material for the sensor was based on considerations of both reliable fabrication techniques and material spectral response. Silicon was selected both for its response from the lower end of :he visible to the long-wavelength cut-off above 11,000A and for its well-developed epitaxial planar-diffused technology. The electrical image is composed of M x N discrete information bits with each bit emanating from one of the individual phototransistor sensors. The interconnected mosaic concept was introduced to provide a structure which is both manageable in its number of leads and compatible with conventional viewing systems which accept XY data. Interconnections along the Y direction are provided by internally diffused strips, while interconnections along the X direction are formed by vacuum-deposited metal surface bars. Unique access to any individual element of the M x N mosaic is available through one of the M and one of the N external leads. PHOTOTRANSISTOR DESIGN A phototransistor structure2 was selected for the discrete photon detectors of which the mosaic is com- posed. Basically, the mode of phototransistor operation is such that incident quanta are absorbed in the base area of the element where they generate electron-hole pairs. If the electrons are liberated within a diffusion length of the depletion region around the collector junction, they will diffuse to this depletion region and be swept across the junction where they account for a small component of the photocurrent. Similarly, the holes which are created within a diffusion length of the emitter junction will diffuse to this region where they forward bias the emitter and so induce injection of minority carriers into the base. This accounts for the substantial transistor photocurrent. This mechanism of operation implies several design and fabrication goals for the individual sensor elements: high electrooptical conversion efficiency, sufficient carrier lifetime in the base regions, and a maximum volume in which photons can be absorbed within a diffusion length from the depletion layer about the junction. Moderately high transistor gain is desirable to permit readout at low light levels; however, very high gain has a deleterious effect on mosaic dynamic range. The mosaic must also be characterized by good uniformity of element response as well as by good isolation between elements. The above set of conditions imply that junction depths must be optimized with respect to photon ab-

Jan 1, 1967

By H. Biloni, B. Chalmers

Metallographic studies of the segregation of solute in and near the surfaces of chill-cast alloys have shown that, under sufficiently severe conditions, nucleation on the cold mold surface is followed by the formation of a disc of solid of the same composition as the liquid. This disc is surrounded by two concentric rings in which the concentration first decreases and then increases. The discs and rings are surrounded by a region in which arms of low solute concentration separate regions of high concentration. These arms start at the periphery of the outer ring and are initially radial; they make a gradual curved transition from the radial direction to a crystallographic direction. A normal dendritic substrcture is established beyond the position at which the arms have become crystallographically aligned. When the cooling is less severe, the disc becomes less irregular in shape and has a concentration minimum at the center. The term "Predendritic" is proposed for these transitions. A qualitative explanation is outlined for the phenomena observed at the surface and in the equiaxed region of the ingots. It is well-established1-' that the normal mode of solidification for alloys under most conditions is dendritic. A characteristic of dendritic solidification is the growth of rods or plates in specific crys-tallographic directions. The remaining liquid solidifies later, usually with a higher concentration of solute than the original liquid. Very little is known, however, about the transition from a crystal nucleus until the dendritic pattern is established. It is clear that there must be a transition because the nucleus is much smaller than the characteristic spacings found in the "steady-state" structure. And it also seems clear that the initial conditions must be important since the undercooling existing during the nucleation of the grains has a remarkable influence upon the grain size4 and the segregation pattern. The present paper reports a study of the micro-segregation in regions where crystals have nucleated, and the conclusions that can be drawn regarding the early stages of growth, in connection with the details of the solidification process. Aluminum of 99.993 purity, A1-Cu alloys containing 0.5 to 20 wt pct Cu, and Pb-Sn alloys with 5 wt pct Sn were cast and examined metallographically at the as-cast surface, close to the surface, and in the interior after casting. The alloys were cast from controlled temperatures against surfaces at different temperatures ranging from room temperature to approximately the liquidus temperature of the alloy. The surfaces were, in various experiments, graphite, stainless steel, copper, and aluminum. The methods used for metallographic examination were as follows. 1) The Nomarski phase contrast interferometer6 was used to obtain information on the topography of the as-cast surface. 2) The distribution of any solute present in A1 99.993 pct and copper in the A1-Cu alloys was studied by the anodic film technique originated by Lacombe and Mouflard7 and previously applied in the study of the solidification of aluminum by Biloni.8 The electropolished surface of the specimen is anodized under controlled conditions so that the thickness of the oxide film and therefore the interference colors under white light are sensitive to the local composition. This technique was applied to the as-cast surfaces and to metallographically prepared surfaces from the interior. 3) If in Al-Cu alloys, the surface is anodically oxidized for several minutes, instead of seconds as in the color-film technique, the copper segregation is outlined by a different level with respect to the matrix. This method is less sensitive but complementary to that for the color films. 4) The distribution of copper revealed and mapped by the metallographic techniques was confirmed by electron probe. 5) In samples with high concentration of copper, evidence for the copper distribution could be seen optically after electropolishing.

Jan 1, 1965