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By L. Guttman, C. S. Barrett, J. S. Bowles

THE transformation from the face-centered cubic (Al) to the face-centered tetragonal (A6) structure in certain alloys of the indium-thallium system reported in the preceding paper1 exhibits many interesting crystallographic and metallo-graphic features, most of which presumably are similar to those of other cubic to tetragonal transformation such as, for example, the one which occurs in the chromium-manganese system,2 the one which, judging by microstructures and X-ray pat-terns,3 probably occurs in the copper-manganese system, the ones accompanying ordering in FePt,4 Copt,&apos; and AuCu,5 and the transformation near 115°C in barium titanate.6,7,8 If, as mentioned in ref. 1, the transformation is second order, the phases in equilibrium do not differ in composition, and it is both possible and likely that the mechanism is a diffusionless one. The present paper reports the results of a detailed study of the indium-thallium transformation and presents a theory for the atomic movements of the transformation which accounts for the observations. Experimental Crystallography of Transformation Markings The transformation from cubic to tetragonal in the indium-thallium alloys produces microstruc-tures of the type illustrated in fig. 1. These structures are observable either as relief effects produced on a smooth surface by the transformation, or after polishing and etching. These microstruc-tures are markedly similar to those found in copper-manganese alloys: iron-platinum alloys," and barium titanate,7,8 where it has been demonstrated that the lamellae are parallel to the (101) planes of the original cubic crystal. That the observed lamellae in indium-thallium alloys are also parallel to the (101) planes was proved by the following analyses performed on single grains in polycrystal-line samples containing 20.75 atomic pct T1. A study of the microstructures of a number of specimens revealed the fact that in many specimens the transformation markings continued unchanged in direction across the (111) interfaces of annealing twins. This at once demonstrates that the lamellae are parallel to a plane that is common to both parent and twin orientations, i.e., a plane in the zone <112>. This conclusion is compatible with the lamellae being parallel to (101) planes. To make a more complete determination of the habit plane of the lamellae, two grains were selected in which annealing twins ({Ill) twins) had been present in the high-temperature modification (cubic). The traces of the annealing twins, and of the transformation markings in both parent and twin crystals, were plotted together in stereographic projection, and a parent-twin pair of orientations was found that accounted for all lamellae as (101) planes. The projection of one of these twinned grains is reproduced in fig. 2. In view of these results the probability of there being another solution for the habit plane remote from (101) is very small, but to explore the possibility that the habit plane deviates slightly from (101), a more precise analysis was undertaken. In this, the orientation of a grain in the cubic modification was determined from a Laue back-reflection photograph taken with the sample at 90°C. The traces of the lamellae that formed in this grain on cooling were plotted in the stereographic projection

Jan 1, 1951

By R. W. Kraft

A comprehensive analysis of the crystallographic and metallographic structure of several Al-CuAl, lamellar eutectic grains solidified under equilibrium conditions has been performed. A consistent pattern was discovered which was interpreted in terms of a detailed atomistic model of an interface describable as: Interface II {111} AL 11 {211} <101>Al <120> ? Several arguments are presented to substantiate the theory. It is compared in detail with another interface model characteristic of the Wid-manstatten precipitate of CuAl2 in A1 at low Cu levels. In addition, a concept is proposed to help explain the structure of the compound CuAl2. It has been shown that the eutectic alloy between A1 and CuA1, can be forced to solidify as substantially parallel lamellae throughout a relatively large volume&apos; by appropriately controlling the solidification parameters of growth rate, thermal gradients, and impurity level and by causing solidification to proceed in a unidirectional fashion from one end of an ingot to the other. A typical ingot solidified in this manner shows three distinct zones along the length of the ingot, Fig. 1. At the head end is a region consisting of many different grains, some of which are elongated slightly in the direction of solidification. After a few centimeters of growth, this polygrained region gradually merges into the second zone which consists of a very few elongated coarse grains. The third zone at the tail end is a region of eutectic colonies caused by segregation of impurities. Microscopic examination of transverse and longitudinal sections of the ingots reveals that during the initial stages of growth (zone I) the lamellar structure is frequently highly distorted and twisted but that it straightens out in zone 11. This characteristic pattern suggests that nucleation of a eutectic grain begins randomly at the head end of the ingot, and that some particular orientation relationship between the phases is more favorable from an energy consideration than others for one reason or another, so that some grains grow and others which are less favorably oriented do not. The favorably oriented or "low energy" grains continue to grow until excess solute rejected at the advancing liquid-solid interface destroys the parallel lamellar structure by forming a cellular rather than a substantially planar interface. When this occurs, the parallel lamellae fan out and eutectic colonies are formed because the lamellae grow perpendicular to the local interface. Questions naturally arise as to the crystallographic nature of the "low energy" grains. What do the grains which persist into zone II have in common? Can this explain the stability of such grains during growth? To answer the first question it is necessary to determine the crystallographic orientation of the lamellae of each phase relative to one another. This has been done for Al-CuAl2 eutectic alloys by two different techniques on specimens prepared in different ways.2&apos;3 The results were different from each other and from a third orientation relationship between A1 and CuA1, reported for overaged A1-Cu alloys.4 The three orientation relationships are summarized below to illustrate their similarities and their differences. (001) ? 11 (001) A1; [l00] ? 11 [l00] Al Eutectic, Ref. 2 (001) 8 1) (001) A1; [1001 ?11 [1101 A1 Eutectic, Ref. 3 (001) ? 11 (001) A1; [l00] ? 11 [310] Al; Widmanstfitten, Ref. 4 It is not surprising that different orientation relationships were found by Ellwood and Bagley,&apos; and Takahashi3 since the solidification conditions were quite different. Nor is it surprising that Mehl, Barrett, and Rhines 4 found still another value since they were studying another type of alloy. Nevertheless, it would be expected that some logical and systematic pattern could be found if all the alloys were at equilibrium since the same two phases are

Jan 1, 1962

By A. E. Dwight

Lattice parameters were determined for eighteen equiatornic alloys of the CsCl-type structure, ten of which were previously un-reported. It was found that fomation of the CsCl-type structure in binary alloys of the transition elements is largely dependent on position of the elements in the periodic table. The relative size of the two elements was not found to be a controlling factor. A recent paper by Beck, Darby, and Arora1 corre-lates the occurrence of CsC1-type ordered structures with the position of the constituent elements in the periodic table for the first long period. It was also suggested that a definite increase in relative bond strength between unlike atoms occurred when, in binary alloys of iron-group elements, the other component is changed from a chromium-group element, to a vanadium-group element, to titanium. A later paper by Philip and Beck2 noted that the lattice contraction increased in the order CrFe, VFe, and TiFe. It was also noted by Philip and Beck2 that the lattice contractions of CsC1-type alloys decreased in the order: TiFe, TiCo, and TiNi, which is an apparent reversal of the contractions expected from the position in the periodic table. It was suggested that the increasing lattice contraction is an indication of increased stability, i.e., greater A-Bbond strength. The present investigation was carried out to determine whether the relation of the position in the periodic table to the formation of the CsC1-type structure was also correct for alloys involving the second and third long-period elements. A systematic search was made for CsC1-type structures among equiatomic alloys and for those found, the lattice contraction was determined. EXPERIMENTAL TECHNIQUE The elements Y, Gd, Ti, Zr, Hf, V, and Cb are designated the A group and the elements Mn, Re, Fe, Ru, Os, Co, Rh, Ir, Ni, Pd, Pt, Cu, Ag, and Au are designated the B group. Equiatomic alloys were prepared for 57 AB combinations. The alloys were arc melted in a multicrucible furnace3 in buttons ranging from 5 to 20 g. Chemical analyses were not made, as the charge weights agreed closely with those of the buttons after melting. The alloy buttons were homogenized at 800° to 900°C. Metal log raphic and X-ray specimens were prepared and heat treated at temperatures from 600° to 1200°C. Specimens for X-ray diffraction were usually ground to a powder in an agate mortar; however, needle-shaped solid specimens were used when the alloy was sufficiently ductile to permit their preparation. Diffraction patterns were taken with a Straumanistype Debye-Scherrer camera using filtered Cu or Co radiation. The lattice parameters were obtained in A by plotting the calculated a0 values against the cos29/sin 0 + cos2?/? function and extrapolating linearly to ?= 900. Metallographic control specimens were polished on cloth wheels with diamond paste and etched with various phosphoric and nitric acid reagents. RESULTS The eighteen equiatomic alloys listed in Table I gave evidence of a cubic structure with two atoms in the unit cell, although two of these cubic structures exist only at elevated temperatures and transform to a tetragonal structure on quenching. Nine of these eighteen alloys gave diffraction patterns with super-lattice lines showing that the structure is of the CsC1-type. The lack of superlattice lines in patterns of the other nine alloys may be attributed to the small difference in atomic scattering power of the components. Metallographic study indicates the occurrence of nine narrow single-phase fields at the AB composition. Any or all of these nine may also have a CsC1-type structure. The VFe alloy was found to have a CsCl-type structure by Philip and Beck2 through the use of CrKa radiation (for which the scattering factor of V is anomalously low), whereas the Cu radiation used

Jan 1, 1960

By T. V. Philip, Paul A. Beck

IN a previous note1 it was pointed out that the available information suggests a distinct correlation between the occurrence of the CsCl-type ordered structures formed in equi-atomic binary alloys of transition elements and the location of the components in the periodic table. A definite increase in the bond strength between unlike atoms, as compared with that between like atoms, is indicated when, in binary alloys of iron group elements, the other component is changed from chromium to vanadium or tantalum, to titanium. Laves and Wallbaum,2 eported a CsCl-type ordered structure for the stable intermediate phase at the equi-atomic composition in the Ti-Fe system (and also in the Ti-Ru and Ti-Os systems). On the other hand, Tagaya, Nenno, and Nishiyama8 found no ordering of this kind in the body-centered-cubic solid solutions in the Cr-Fe system. It was concluded&apos; that the occurrence of a CsC1-type ordered phase of intermediate stability may be expected in the V-Fe system. Preliminary results were described: suggesting that at least some degree of order does exist in the equi-atomic alloy VFe, annealed at 1200° C and quenched in cold water. Using CrK radiation, at least one weak superlattice line was found in the X-ray diffraction pattern which could be indexed as (100). and which disappeared, as expected, when FeK radiation was used. It was found more recently that the impurity phase (presumably VN), previously giving rise to additional X-ray diffraction lines and thus preventing definite indentification of the (111) superlattice line,&apos; can be eliminated by preparing the alloy by arc melting in a helium atmosphere. The alloy specimen melted in this manner had clean surfaces and showed very littte, if any, nitride phase in the microstructure after annealing in vacuum at 1250°C and quenching. The relative intensity of the (100). superlattice line with CrK radiation was greatly increased and, in addition, a fairly strong (Ill) , superlattice line and even a detectable (210) , super- lattice line was obtained, when the powder produced by filing from the 1250°C annealed and quenched specimen was reannealed in vacuum at 625°C for 1 hr. By using a cellulose acetate film as a filter in front of the photographic film in order to absorb the soft radiation resulting from air fluorescence, the background intensity could be reduced considerably in relation to the intensity of the diffraction lines. All three superlattice lines, (100),, (111)., and (210)a, disappeared, as expected, when a radiation other than CrK, was used.1 For the powder reannealed at 625ºC, the strongest lines of the s pattern also appeared weakly, indicating that during an: nealing at 625°C a small amount of transformation took place from the body-centered-cubic to the u phase. However, since all superlattice lines to be expected in the 0 angle range used were actually very clearly present, it was evident that the high temperature body-centered-cubic phase had transformed into a CsCl-type ordered structure before much s phase could form. Reannealing filings for 1 hr at 700°C resulted in complete transformation into the u phase. After 1 hr at 650°C the transformation into IT was not quite complete, the strongest body-centered-cubic diffraction lines were still present, although fairly weak, and only the (100), super-lattice line was detectable (extremely weak). These findings not only confirm the previous conclusionsL as to the existence of a CsCl-type ordered structure in the Fe-V system, but they also indicate that this metastable ordered structure forms at about 600°C preceding the formation of the stable s phase, while at 1250°C the disordered body-centered-cubic phase is stable. The faint (loo) , super-lattice line previously observed in specimens quenched from 1200°C was probably due to the fact that the quenching of the powder in an evacuated fuse:! quartz capsule was not fast enough to prevent completely the ordering reaction on cooling. The present results conform with the suggestion that the A-B bond, in relation to the A-A and B-B bonds, increases in strength from CrFe to VFe to TiFc. It is, therefore, of interest to compare the measured values of the average atomic diameter for these three alloys with the corresponding average atomic diameters calculated from the atomic radii of the elements concerned. ASSUming A-B contacts

Jan 1, 1958

By Jean Howard, E. V. Walker

ManY papers have been published during the last few years on the formation of cube texture in 3 1/4 pct Si-Fe, e.g., those of Assmus, Detert, Dunn, and Walter.1,5 All are concerned with the formation of the texture by a process involving secondary re-crystallization and the relative surface energies of (100) and other planes under certain conditions of annealing. The authors have found that cube texture can also be produced by heavily cold rolling strip in which the cube-on-edge texture has first been formed. Cube texture develops on annealing, but in this case it is not accompanied by secondary recrystallization. The starting material is a sintered compact, the preparation of which has been described elsewhere. Using a conventional processing schedule employed for producing commercial silicon-iron with the cube-on-edge texture, the sintered compact is cold rolled to such a thickness that, after annealing in pure hydrogen for 24 hr at 1350&apos;~ to produce the cube-on-edge texture, the final thickness required can be obtained by giving a further cold reduction of 85 pct. Cube texture is developed by annealing at -900~~ in pure hydrogen: it has been produced in strip up to 0.024 in. thick. The microstructure of the strip consists of a mass of small primary grains within the boundaries of the much larger secondary grains which ex- hibited cube-on-edge texture before the final cold roll. If the strip is annealed at -1000 after the 85pct reduction, secondary recrystallization begins to take place, provided a strong cube texture is present. On raising the temperature more secondaries grow, complete secondary recrystallization being obtained on annealing at -1150&apos;. From this it is concluded that the secondary crystals are produced by the texture inhibition mechanism: they show (110) and (111) planes in the surface of the strip, the former being the more numerous. It has been observed that the more perfect is the (110) [001] texture prior to cold rolling by 85 pct, the further from the ideal is the (100) [oo].] texture. A similar treatment applied to high quality commercial grain-oriented silicon-iron strip failed to produce cube texture. The tilt and rotation of the [loo] from the rolling direction was less than 10deg, but (to within a few degrees) the planes (130), (450) and (210) lay in the surface of the strip. Acknowledgment is made to the Engineer-in-Chief of the British Post Office for permission to use the information contained in this note.

Jan 1, 1962

By S. R. Goodman, Hsun Hu

PREVIOUS investigations on the temperature dependence of rolling texture transition in high-purity silver1&apos;2 and in electrolytic copperS have shown that a brass-type texture is favored at low rolling tem- peratures, whereas a copper-type texture is favored at high rolling temperatures. The transition takes place gradually as a function of the temperature of deformation, and a general correlation between texture transition and stacking-fault frequency or probability exists—a copper-type texture is associated with low stacking-fault frequencies, and a brass-type texture is associated with high stacking-fault frequencies. If the stacking-fault energy is calculated from the stacking-fault frequency and the dislocation density of the specimen, one finds that with low stacking-fault energies a brass-type texture will be developed, and that with high stacking-fault energies the texture will be of the copper-type. These findings are consistent with the idea that the formation of rolling textures in fcc metals may depend on their stacking-fault energy, although results of such nature always involve the ambiguity that a change in the density (or width) of stacking faults produced by deformation at different temperatures may be actually due to a variation in the stress, rather than due to a change in the specific stacking-fault energy per se. It has also been shown that the type of annealing texture depends entirely on the deformation texture. From a copper-type rolling texture, a strong cube texture will be formed upon recrystallization, whereas from a brass-type rolling texture, the annealing texture will be quite different from the cube orientation. Based on these results on texture transition in pure fcc metals, it appeared possible to effect a

Jan 1, 1963

By Y. C. Liu

IN the study of the rolling recrystallization of face-centered-cubic metals, a most unusual observation is the frequent occurrence of a sharp re-crystallization texture of cube orientation. Although the general conditions under which this texture occurs are known,2,3 its origin is still not well understood. Based on the observed relationship between the {110}<112> deformation texture and the (113}<211> recrystallization texture, regardless of the effect of alloying elements2,4,5 or the symmetry of the deformation texture,&apos; there is a strong indication that a specific component of the deformation texture is related to a specific component of the recrystallization texture. Using the rotational reorien-tation relationship as a means of correlating the deformation and recrystallization textures, it was proposed that the deformation component which related the cube texture has {358)<523> indices."&apos; The investigation reported here was undertaken to seek experimental evidence for the validity of this proposition. Experimental Procedure Two copper single crystals having an initial orientation within 2" of (358) [523] indices were prepared in the manner described in Ref. 5. The purity of the copper was between 99.94 and 99.97 pct. The original dimensions of the crystals were as follows: specimen A, 0.465 in. thickness, 0.844 in. width, and 0.955 in. length; specimen B, 0.535 in. thickness, 1.120 in. width, and 1.36 in. length: Rolling was done on a laboratory rolling mill according to the following schedule: 0.010 in. per pass to 50 pct reduction in thickness, 0.005 in. per pass to 70 pct reduction inthickness, 0.002 in. per pass to 85 pct reduction in

Jan 1, 1958

By P. K. Koh

With identical annealing heal treatment the development of major annealed texture component seems to depend primarily upon the degree of cold reduction. Cube texture was evident on annealing- 1/2-, 1,/4- , and 1/8-mil tapes at 760°C, but it dissap-)eared on 927°C annealing of 1/8-mil tape cold reduced 96.9 pct. Some cube texturee remained as tile strongest 927 °C annealed texture component in 1/4- and 1/2-mil tapes cold reduced 93.8 and 87.5 pet, respectively. The switching coefficient increases and the squareness ratio of the hysteresis loop decreases proportionally with increasing cube pole concentration. PREVIOUS studies 1,2 on the 1/8-mil ultra-thin molybdenum permalloy tape established cold rolled and annealed textures and their magnetic performance for switch-core and coincident-current memory applications, The major cold-rolled texture was found to be of {110}<335> which is quite similar to the cold-rolled texture of silver and 70-30 cartridge brass. The major annealed texture components were found to be of {120)-- <001> and {113)<785>. Although the direction for easy magnetization for this type of alloy by the single-crystal technique is not avaliable, knowledge from related studies indicated [Ill] as the easy magnetization direction, similar to that of nickel.&apos; Therefore, a cube orientation, if present in the tape core, would be most detrimental to its magnetic performance.3 In its cold-rolled state the tape had < 111 > directions in the rolling plane 14.5 deg from the rolling direction.&apos; But the coercive force of the tape in such condition will be excessive due to large amounts of internal stresses, and therefore makes its use impossible for memory applications due to long switching time. In lower nickel alloys, suchas 50-50 nickel-iron and 65 pct Ni-Fe sheets, previous investigators4"7 found the presence of cube texture on annealing at 900° to 1000°C. Although no cube texture was found in annealed 1/8-mil tape,"&apos; its presence in some annealed 1,/4-mil tape cores was firmly established a year ago as verified by significant differences in magnetic properties. The present investigation is devoted to the orientation study and related magnetic measurements on three thicknesses of the ultra-thin tape, namely, 1/8 mil, 1/4 mil, and 1/2 mil, cold rolled directly from annealed 4-mil stock in the cold-rolled state as well as after they were annealed at 760°C (1400° F) and 927°C (1700°F). The results obtained might serve to elucidate the development of annealed texture as well as to establish clearly the importance of preferred orientation on the magnetic performance of such memory tape core despite its well-known low magnetic anisotropy exchange energy. EXPERIMENTAL PROCEDURE Samples were taken from one production heat (0.010 wt pct C, 0.004 P. 0.016 S, 0.80 Mn. 0.29 Si. 79.69 Ni. 4.42 Mo. 0.054 Co, 0.002 Ti, 0.12 Al, 0.04 Mg, 14.56 Fe). Annealed 4-mil sheets were cold re-

Jan 1, 1962

By P. A. Beck, M. V. Nevitt

All binary and a number of ternary u alloys formed by first long period transition elements were examined and found to be ferromagnetic at low temperatures. The Curie temperatures for these alloys were measured. It was found that the correlation of the Curie temperature with composition may be described qualitatively in terms of the contributions of the various component atomic moments to the saturation moment of the alloy. Ternary Fe-Cr-Mo u phase alloys were also found to be ferromagnetic, with a Curie temperature dependent on the iron content and largely independent of the unit cell dimensions. Similar relations were found for the Co-Fe-Cr and the Co-Cr-Mo a alloys, where the Co + Fe content appears to determine the Curie temperature. RECENTLY, it has become increasingly evident that the u phase is some kind of an "electron compound." Several papers1-&apos; discussed the empirical correlation between the composition of the various u phases in alloys of the transition elements and the distribution of electrons in the partially filled d-levels of the component atoms. However, a picture of the latter can be formed only in a more or less intuitive manner, as in the Pauling theory," and the rigorous treatment of these problems is at present prevented by lack of adequate knowledge of the electronic structure of the transition elements in the metallic state. It appeared reasonable to expect that investigation of the magnetic properties, which are related to d-shell electrons, may help in clarifying the physical nature of the foregoing empirical correlation. Qualitative tests of samples of various compositions have shown that certain u alloys become ferromagnetic at liquid nitrogen temperature." The present investigation was undertaken to determine whether or not low temperature ferromagnetism is characteristic of a large variety of u alloys and to establish how the Curie temperatures of ferromagnetic u alloys are influenced by changes in composition and in lattice parameter. In the first portion of the in- vestigation, Curie temperatures were measured for a series of u alloys containing either vanadium or chromium with one or two of the elements manganese, iron, cobalt, or nickel. In the second phase of the work, data were obtained on the Curie temperature of ternary (Fe, Co, Cr)a, (Fe, Cr, Mo)a, and (Co, Cr, MO)U alloys, as a function of composition and of interatomic spacing. Experimental Procedures High frequency vacuum melting in high purity re-crystallized alumina crucibles was employed in the preparation of the u alloys, using metals whose lot analyses are given in Table I. Approximately two-thirds of the alloys were analyzed chemically, the analytical results being in acceptable agreement with the intended compositions, as seen in Tables I1 and 111. The specimens were annealed for homogen-ization in an atmosphere of purified 92 pct He and 8 pct H, mixture or in evacuated and sealed Vycor tubes at temperatures from 780" to 1300°C, depending on the temperature range of the u phase field for the particular composition. The annealing temwerature and time for kach alloy are given in Tables 11 and 111. Curie points were determined by an induction method; the change with temperature of mutual inductance between two coils was observed when the i-u specimen was inserted into the flux path between the coils. The transformer consisted of three coils each containing 1200 turns of 0.0050 in. diam enameled copper wire. The center coil, P in Fig. 1, served as the primary coil, and was linked magnetically with each of the two secondary coils, S, and S? which were connected in series opposition. The primary coil, P, was energized by an audio-oscillator, set at a frequency of 500 cycles per sec, while the signal from the transformer secondary circuit,

Jan 1, 1956

By L. F. Coffin

Monotonic tension and controlled cyclic-strain tests have been conducted on an annealed steel at various strain rates and temperatures in the weep range. Evidence of a strain-induced precipitate reaction abozle 500°C was determined from pronounced cyclic-hardening effects, increase inflow stress, inter granular fracture, drastic decrease in low-cycle fatigue resistance, and loss in tensile fracture ductility, as well as through electron-microscopy examination. It was found that strain rate as well as temperature had a pronounced effect on these observations. A parameter of the form P = T(7.95 - log ip) was determined, which was directly related to the ductility, defined either as the true strain at fracture in the case of simple tension, or as 2N 1/2 ? E p in the case of cyclic plastic strain. A highly satisfactory comparison of the cyclic life as measured and as computed from the ductility predicted by this parameter was obtained for a range of temperatures from 400° to 650°C and for crosshead speeds of 0.002 to 0.2 in. per min. MANY metals subjected to sustained loading at elevated temperatures derive resistance to creep from solid-state reactions of various alloying constituents. These reactions can be many and varied, and occur as a consequence of time and temperature according to reaction-rate principles. Strong evidence exists that many of these reactions are influenced by plastic strain, the argument being put forth that dislocations act as sites for nucleation and growth for the new phase.&apos; Many examples can be cited where plastic strain during transformation influences the structure, and as a consequence affects the properties.2 The low-alloy steels serve as particularly good examples for investigating the effect which aging reactions have on resulting properties. Alloying elements such as carbon and nitrogen, and carbide formers such as manganese, chromium, molybdenum, and so forth, can, by the particular aging reaction which occurs, produce strengthening effects at various temperatures, and so endow strength to the alloy over a broad range of temperatures. These effects have been systematically studied by Glen both during creep3 and for constant strain-rate tension testing.4,5 The work of Glen is of considerable interest in connection with the present study, particularly in regard to the effect of various alloying additions on the true stress-strain curve and on fracture ductility at various elevated temperatures. He has shown that strain-age hardening occurred during plastic straining, which was quite specific both to the temperature of the test and to the alloying element involved. When strain-age hardening was noted at a particular temperature, a .minimum in reduction of area was also obtained. Provided the temperature for the minimum reduction of area exceeded 300°C, all minimum ductility fractures were found to be inter-crystalline in nature. For carbon steels containing molybdenum, two minima in reduction of area were noted, one occurring at 200° to 250°C, the precise temperature depending on the percentage of molyb-

Jan 1, 1964

By C. E. Feltner

In a recent note, Meleka and Hiarris1 have reported that the application of a constant strain rate Ad a superimposed cyclic stress to 99.99 pct polycrystal-line zinc at room temperature produced some marked effects in the resultant behavior. In their work, one end of the specimen was pulled in tension at a constant strain rate while the other end was subjected to a sinusoidally varying load (20 cps) of constant amplitude. For the slower strain rates and the larger superimposed cyclic loads, it was observed that the specimen could withstand strains prior to failure more than 100 pct greater than the fracture strains obtainable in simple unidirectional straining. This ductility effect was interpreted as being due to a cyclic stress enhanced thermal recovery process. It is the purpose of the present note to describe some related effects in 99.99 pct polycrystalline cadmium. Specimens were loaded in repeated tension at room temperature with a stress of sufficient magnitude to cause failure in about 1000 cycles and the accumulated strain was measured as a function of cycles and time. In addition, some specimens were subjected to a steady creep stress equal in magnitude to the peak value of the cyclic tensile stress. The resultant strain-cycle and strain-time curves for both types of test are shown in Fig. 1. It should be noted that the static creep specimen failed in about one third the time required to produce a cyclic failure but in the cyclic test the specimen was able to withstand a uniform strain before necking more than 50 pct greater than that possible

Jan 1, 1963

By Randolph H. Schnitzel

MARINGER and schwopel have investigated the internal friction of high-purity molybdenum. After light tensile deformation, they observed internal-friction peaks at about 230° to 290°C, 310" to 370°C, and 820° to 870°C, which they have attributed to the interstitials carbon and nitrogen, and to stress relaxation across grain boundaries, respectively. The electron-beam melting technique has been successfully applied to refractory metals including molybdenum. By this technique single crystals and higher-purity levels than those of commercial molybdenum are readily obtained. Since electron-beam processed molybdenum possesses such desirable qualities, it seems reasonable to determine: u) if internal-friction peaks do exist in these single crystals, and b) if so, whether or not these peaks can possibly be attributed to the motion of specific interstitials. The molybdenum crystal used in this investigation received four melting passes in the electron-beam furnace. This crystal was about 11 in. long and had an orientation 4 deg from the [loo]. The process of making wire consisted of grinding the crystal rod to a diameter of 0.040 in. and then electropolishing to a final diameter of 0.030 in. Since Maringer and schwopel have indicated that the peak at 230" to 290°C may be due to the presence of carbon, the wire was cut into two equal lengths and one length was carburized prior to the internal-friction measurements. The critical condition necessary for the controlled addition of carbon proved to be a difficult problem. Therefore, a simpler nonequilibrium method was used. This consisted of carburizing at 2000°C in graphite powder to obtain a case, annealing at higher temperatures in vacuum to partially dissolve carbon into the molybdenum, and then removing the remaining case by dissolving it elec-trolytically. After this treatment, metallographic examination of a small section of the specimen revealed carbides in the matrix. Vacuum fusion and carbon analyses for both the virgin (as melted) and carburized specimens are given in Table I. Internal-friction measurements of these specimens were conducted in an apparatus of conventional

Jan 1, 1964

By A. S. Yue

The interface morphology of the Mg-32 wt pct A1 eutectic has been studied in terms of the freezing rate, the temperature gradient, and the impurity content. For the impure eutectic it has been found that for a given temperature gradient a fast freezing rate favors the formation of a cellular intwface with sub-grain strucfure. This cellular interface can be changed to a planar interface of the same substructure either by decreasing the freezing rate or by increasing the temperature gradient. For the eutectic of highest purity, as purified by zme-refining, no cellular interface was found for any growth condition and only a planar interface with subgrain structure was observed. Furthermore, a layer of liquid adhering to the interface during de-cantation freezes with shapes and sizes of the phase particles different from those in the main eutectic solid. WHEN a molten metal or a binary eutectic liquid with zero freezing range is solidified unidirec-tionally, a cellular interface is usually formed-3 This interface is schematically represented by Fig. 1. The profile of the solid-liquid interface has a corrugated shape, as shown in Fig. l(a), while the decanted interface consists of cells, as shown in Fig. l(b). The cell boundary in a frozen eutectic is distinguishable from the cell interior by changes in particle shapes, as shown in Fig. 4. The formation of a cellular interface, according to Weart and ack,&apos; is attributed to a layer of constitutionally supercooled" liquid generated at the advancing interface by the simultaneous rejection of an impurity by the two phases of the eutectic mixture, which then permits freezing over an appreciable temperature range. This concept has been further elaborated by Tiller, who predicts that even in the absence of an impurity the rejection of solute ahead of each lamella of the eutectic may promote the formation of a cellular interface. More recently, Chilton and Winegard, and

Jan 1, 1963

By R. M. Hudson

A study has been made of the decarburization of low-carbon sheet steels at temperatures from 1100" to 1475°F in mixtures containing hydrogen, nitrogen, water vapor, and in some instances carbon monoxide. Decarburization occurs principally by the reaction C (in steel) + H20 = CO + H2 and follows a first-order rate law. The decarburization rate appreciably increases with an increase in temperature up to about 1325°F, but the increase in rate is small with fur- Use of a dry hydrogen atmosphere results in a relatively slow decarburization of steel and use of a wet hydrogen atmosphere results in relatively rapid decarburization. It is now generally recognized&apos; that the reaction C (in steel) + H2O = H2 + CO [l] accounts for most of the decarburization that occurs in wet hydrogen-containing atmospheres. The present article is concerned with the decarburization of low-carbon sheet steel in H2-N2 atmospheres ther increase in temperature. The limiting rate of decarburization can be predicted from a knowledge of the solid solubility and diffusivity of carbon and this rate is approached as the water-vapor content of an atmosphere is increased. Hydrogen contents above about 8 pct are beneficial only in that they permit higher levels of water vapor to be used without oxidation of the steel. A buildup of CO in the atmosphere lowers the rate of decarburization. containing various amounts of water vapor and carbon monoxide. The effectiveness of these atmospheres and the influence of temperature on decarburization rate will be considered. EXPERIMENTAL WORK The chemical compositions of the three steels used in decarburization-rate studies are listed in Table I. Steel A was a commercially hot-rolled rimmed steel that was cold-reduced at the Laboratory to furnish specimens 0.025, 0.041, and 0.059 in. thick. Most of the experiments in which this steel was used were made on the 0.041-in.-thick material. Steel B and Steel C were cut from 0.010-in.-thick commercially cold-reduced coils that were

Jan 1, 1963

By A. S. Nowick, A. E. Roswell

Anelastic measurements of atomic mobility in an Ag-Zn substitutional solid solution, make possible a study of the rate of decay of lattice defects frozen into the alloy by quenching. A WIDE variety of experimental evidence has accumulated recently in support of the belief that atomic mobility or diffusion in face-centered cubic lattices occurs principally through the movement of lattice defects, as contrasted to a mechanism by which two or more atoms directly exchange lattice positions. The most direct evidence, summarized in a recent review,&apos; includes the Kirkendall-effect measurements and anelastic measurements on quenched alloys. Theoretical calculations&apos; have focused attention on the lattice vacancy as the defect probably responsible for diffusion in the common metals. Later calculations3 have also shown the possible importance of combined pairs of vacancies in producing diffusion. The only defect that has been considered seriously, in addition to the vacancy and combination of vacancies, is the "interstitialcy" (an atom which has left its normal lattice position and entered an interstitial position) ; the calculated energy of formation of an interstitialcy in a close-packed lattice is too large to permit a reasonable probability for the occurrence of this defect In view of these calculations it may be tentatively concluded that the defects producing atomic movements are single or multiple vacancies. Anelastic measurements of atomic mobility in substitutional solid solutions provide a valuable method for the study of lattice defects. By means of this method it was first demonstrated&apos; that defects may be frozen into a lattice by rapid quenching from high temperatures. (It was subsequently shown6 that resistivity measurements may also be capable of detecting the presence of frozen-in defects.) A summary of the principles involved in these anelastic measurements will now be presented. The anelastic method is based on a relaxation phenomenon in substitutional solid solutions, first discovered by Zenera in a brass and subseauentlv investigated extensively in a series of Ag-zn solid solutions.&apos; This relaxation effect may be observed in the usual way either dynamically, as a peak in

Jan 1, 1954

By F. R. Brotzen, A. R. Troiano, E. L. Harmon

Precipitation processes leading to drastic property changes are a frequent occurrence in titanium alloys containing large amounts of the retained high temperature P phase. In order to establish the kinetics and nature of these precipitation processes, a series of Ti-V alloys was studied. Time-temperature-transformation diqgrams were obtained for 12.5 and 15 pct V-Ti alloys. These diagrams consisted of two intersecting C-shaped curves: one in the high temperature region for the euilibrium transformation P +u+P, and another at lower temperatures depicting the formation of the meta-stable transition phase, o. Presence of the transition phase retarded the precipitation of the equilibrium a phase. Deformation of retained P resulted in the formation of the o phase. Reversion and retrogression phenomena of the transition phase were observed. A negative temperature coefficient of resistance was found in quenched alloys containing from 15 to 20 pct V. CERTAIN titanium-base alloys can be heat treated to produce appreciable changes in properties. For example, remarkably high hardness values may be obtained by aging or isothermal quenching treatments. Most of the information available to date on the hardening of titanium-base alloys has been systematized recently in an article by L. D. Jaffe.&apos; It was pointed out by that author that the hardening was the result of the decomposition of the retained j3 phase to the equilibrium constituents, a and p, and that this decomposition had many of the characteristics of an age-hardening reaction. Recent work by Frost et al.&apos; 3 has contributed much to the fundamental understanding of the hardening mechanism in titanium-base systems. These authors established the presence of a transition phase in the transformation from j3 to a + j3, which they designated o. The transition phase, in their opinion, was at least partially responsible for the hardening and embrittling effects which accompanied low temperature aging. The present investigation of the mechanism and kinetics of the hardening reaction was carried out for the Ti-V system. Ti-V alloys were ideally suited for a study of the decomposition of j3, since age hardening was known to occur readily in this system." Moreover, no intermetallic compounds occur in this system, Fig. 1, thus simplifying the interpretation of the results. Material and Procedure Alloys containing 12.5 and 15 pct V were prepared by Battelle Memorial Institute from iodide titanium furnished by Watertown Arsenal and vanadium of 99.8 pct purity. Ingots of 8 Ib each were prepared by arc melting in an inert atmosphere with a tungsten electrode and water-cooled copper crucible. Each of the ingots was melted once, forged at 1600°, and rolled at 1400°F. Following this, the ingots were remelted, forged, and then rolled to % in. sq bars, which were subsequently grit blasted, pickled, and wire brushed. Another set of alloys containing nominally 17.5 and 20 pct V was produced by the Research Div., New York University. These alloys were arc melted in an inert atmosphere in lots of 20 g each. Heat-Treating Procedure: Whenever feasible, specimens were heat treated in high vacuum resistance furnaces at pressures below 5x10." mm of Hg. Aging and isothermal quenching treatments were generally carried out in salt baths. Specimens were not attacked by the salt, provided the tempera-

Jan 1, 1956

By J. E. Harris, J. A. Whiteman, A. G. Quarrell

The graphitizing reaction has been studied in a number of Fe-C-Si alloys in the temperature range 550° to 725°C. The TTT relationships exhibit "C" curve behavior except where nuclei, either foreign or graphite itself, are present before the start of the reaction. It is suggested that the rate -controlling process is diffusion of carbon through ferrite. For the reaction to proceed at the observed rate it is calculated that the matrix surrounding the particles must deform plastically; experimental evidence for this deformation is presented. It is well-established that the decomposition of austenite to pear lite occurs at a maximum rate at an intermediate temperature.l-5 This "C curve" be- havior, which has been reported in a number of other systems,5-10 is a consequence of the relatively low rate of nucleation of the pearlite at small degrees of undercooling of the austenite. The cemen-tite phase in the pearlite is, in turn, metastable with respect to graphite and if steels are held for prolonged periods at elevated temperatures decomposition of the pearlite occurs. It is difficult to predict the shape of the subcriti-cal graphitization TTT (time, temperature, transformation) diagram for, in this case, undercooling does not occur. However, the free-energy change for the reaction does decrease with increasing temperature" (Fig. 1) and it is thus anticipated that the rate of nucleation will fall as the temperature approaches that of the eutectoid. In addition, a decreasing growth rate with increasing temperature is predicted in some circumstances with pure Fe-C alloys. However, apparently similar conditions exist for first-stage graphitization in the austenite range (malleabilization) but no evidence for C curve behavior has been presented in the extensive liter-

Jan 1, 1965

By J. D. Mote, J. E. Dorn

This investigation was undertaken to study the effects of piledup arrays of dislocations on inducing slip, twinning, and fracturing in magnesium bicrystals. A series of variously oriented bicrystals of magnesium having a vertical grain boundary were prepared and tested in tension. It was found that piled-up arrays of dislocations at the grain boundary could, under appropriate conditions, induce slip, twinning- and cracking. The results that were obtained substantiate, at least qualitatively, the general dislocation mechanism for transmission of strain across grain boundaries and the Petch-Stroh concept of fracturing. WHEREAS single crystals of magnesium generally exhibit extensive deformation, coarse-grained poly-crystalline magnesium at subatmospheric temperatures fractures after a few percent elongation.&apos; Although a small amount of ductility is obtained, several features of this fracturing are characteristic of typical brittle behavior. Over a rather broad temperature range the fracture stress is insensitive to the test temperature and the fracture stress increases linearly with the reciprocal of the square root of the mean grain diameter. The course of fracturing is predominantly intergranular, but small fragments of adjacent grains frequently adhere to the fractured surface.2 The brittle behavior of polycrystalline magnesium is attributable to the limited number of facile deformation mechanisms it exhibits at low temperatures. For a general deformation of a randomly oriented polycrystalline aggregate, each grain must exhibit at least five independent mechanisms of deformation to permit accommodation of the imposed deformation from grain to grain.= Although minor amounts of prismatic slip occur in corners of grains where stress concentrations are known to be high, glide in polycrystalline magnesium at low temperatures takes place almost exclusively by basal slip.&apos; The common type of twinning, which takes place on the (1012) pyramidal planes, can under the most favorable orientations, lead to a. strain of only 6.9 pet; the contribution of twinning to the tensile strain would indeed be much less than this in a randomly oriented polycrystalline aggregate of magnesium. Since the three mechanisms of basal slip are coplanar, they are equivalent to only two independent mechanisms, a number insufficient for a general deformation. Consequently, once the permissible twinning has taken place in conjunction with basal slip, no further plastic deformation is possible because of interference to slip at the boundaries of dissimilarly oriented grains. At this stage brittle fracturing takes place due to high stress concentrations at the juncture of slip bands with the grain boundaries; the predominance of intergranular fracturing in magnesium, in preference to transcrystalline fracturing which is prevalent in zinc, has not yet been rationalized. A more atomistic description of the plastic behavior and fracture characteristics of magnesium follows from the analyses made by stroh4 on the stresses induced by piledup arrays of dislocations. Slip first takes place by dislocation motion in the most favorably oriented grains. As the dislocations approach the boundary of a dissimilarly oriented adjacent grain they begin to form an array of dislocations with its attendant stress field. Piledup arrays of screw and edge dislocations introduce high localized shear stresses at the spur of the array; piledup arrays of edge dislocations also induce high tensile stresses localized in the vicinity of the grain boundary. Whereas the shear stresses can induce slip to take place, the tensile stresses, if sufficiently high, can cause fracturing. The localized shear stress will be relieved if sufficient numbers of mechanisms of deformation become operative in the original and the adjacent grain to permit accommodation of the dislocations in the grain boundary. In this event a ductile behavior will be obtained. But if the number of deformation mechanisms is insufficient for complete migration of dislocation arrays into the grain boundary, the tensile stresses due to the edge components of piledup dislocation arrays will continue to increase with increasing applied stress until fracturing takes place. Whereas face-centered-cubic metals have a sufficient number of mechanisms of slip for accommodation of dislocations in their grain boundaries to exhibit ductile behavior, hexagonal-close-packed metals, in general, do not. Consequently, hexagonal-close-packed metals are usually brittle except when conditions such as alloying or temperature permit facile slip by a number of mechanisms. The arguments presented above suggest that the mechanical behavior of magnesium depends on whether or not dislocation arrays in adjacent grains can enter the grain boundary. When such accommodation is possible, ductile behavior is expected; but when such accommodation is impossible, fracturing will ensue. To further test the validity of these arguments it was considered advisable to study the

Jan 1, 1961

By N. S. Stoloff, M. Gensamer

The effects of temperature, grain size, and magnesium content on the strength and ductility of cadmium were studied in the range -269° to 23 °C. A sharp drop in ductility between -140° and -190°C marked a transition in fracture mode from ductile shear to inter gvanular fracture. The addition of up to 15.35 wt pct Mg in solid solution raised the transition temperature, and for two compositions produced basal cleavages. The relatively high ductility of unalloyed cadmium at cryogenic temperatures is attributed to temperature independence of the yield stress and the absence of dislocation locking. ALTHOUGH numerous studies have been made of the mechanical properties of cadmium single crystals,1-5 no comprehensive investigation of deformation and fracture in polycrystalline cadmium has been carried out. This investigation was initiated to determine the mechanisms of deformation and fracture in polycrystalline cadmium, and to attempt to correlate these mechanisms with the behavior of other metals of hexagonal close-packed structure. A second objective of this work was to determine whether the slip systems and mode of fracture of polycrystalline cadmium could be altered by reduction of the axial ratio, which is 1.886, through alloying. Dorn and co-workers&apos; have been able to induce prismatic slip in magnesium through alloying to reduce the axial ratio. Single crystals of cadmium are so highly ductile that glide strains of 400 pct are observed at room temperature, and 80 pct at -269°c.l Failure in every reported case is by necking down and shearing off of the crystal. Cleavage in cadmium single crystals has never been observed. Magnusson and aldwin reported a ductile to brittle transition for polycrystalline cadmium at about —160 C for specimens tested at a strain rate of 0.05 in. per in. per min in tension. However, even at -196° substantial ductility was retained. The transition temperature was strain rate dependent, impact loading at 19,000 in. per in. per min raising the transition temperature to about -85°C while leaving the low temperature ductility unchanged. No information was given on microstructures or fracture mode, nor were stress-strain characteristics reported. In the present investigation, tensile tests were carried out on polycrystalline cadmium and solid solution Cd-Mg alloys to determine flow stresses, work-hardening behavior, fracture strength, and fracture ductility. Variables employed included temperature, alloy content (and c/a ratio), and grain size.Some compression data also have been obtained for comparison purposes. I) EXPERIMENTAL PROCEDURES A) Polycrystalline Cadmium. Two purities of cadmium were employed in this investigation. Cadmium of 99.95 pct purity was supplied by Belmont Smelting and Refining Co. Cadmium of 99.99+ pct purity was supplied by American Smelting and Refining Co. Threaded cylindrical tensile specimens were machined to give approximately a 1/2 in. long straight gage section with a reduced diameter of about 0.110 in. Specimens were heat treated in evacuated pyrex capsules in the range 140" to 290°C to produce desired grain sizes. All specimens were polished in a solution containing 320 g Cr03 and 20 g NaS04 in 1000 cc of H20. Occasionally this solution caused staining of the specimen, in which case a further polish with a solution of one part HN03 and two parts H202 in two parts ethyl alcohol was employed.&apos; Tensile testing was carried out on a screw-driven Instron machine, at a constant rate of crosshead movement, which for most tests was 0.01 in. per in. per min. Load vs time, and therefore load vs total strain, were recorded. Temperatures from -196" to -40°C were obtained by passing nitrogen gas through a copper coil immersed in liquid nitrogen, and then feeding the gas into a jacket around the specimen. Tests at -269°C were carried out in a helium cryostat, at a crosshead rate of 0.1 in. per in. per min, with specimens of 1 in. gage length. Compression specimens were tested in a jig similar to one designed by ilman. The jig was mounted on the under side of the Instron crosshead. Load was transmitted by means of a rod connected to a movable part of the jig at one end, and the load cell at the other end. A downward movement of the crosshead applied a compressive load to the specimen.

Jan 1, 1963

By Robert Corcoran, George Wiener

A study of the orientations and microstructure of 3 pct Si Fe alloys after deformation and recrystallization has been made. The components found after deformation agreed with recently published work on single crystals, being primarily (001) [110], (111) [110], and (111) {112}. A pronounced effect of temperature of recrystalliza-tion was found on microstructure and orientation. Low temperature recrystallization was interpreted in terms of low angle boundaries which resulted in the retention of a large amount of the deformation texture. High temperature recrystallization was interpreted in terms of high angle boundaries which gave rise to new components. No simple angular relation between the deformation and high temperature recrystalliza-tion textures could be detected. DEFORMATION and recrystallization behavior of 3 pct Si-Fe has been the subject of scientific investigation for many years. The fact that it is a single phase body-centered-cubic alloy from room temperature to the melting point permits study over a wide range of temperature without the complications of phase transformations. There is an important commercial interest in this material as well, since it is widely used in the electrical industry. In the present work the deformation and recrystallization behavior of poly crystal line silicon iron has been studied with the primary purpose of evaluating the effect of a second phase impurity on both processes as well as the effect of temperature on the recrystallization characteristics. With regard to deformation, the emphasis has been placed on a study of the crys-

Jan 1, 1957