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## Material Information- Title:
- Guinier-Preston zone evolution in 7075 aluminum
- Creator:
- Healey, John Thomas, 1950-
- Publication Date:
- 1976
- Copyright Date:
- 1976
- Language:
- English
- Physical Description:
- xii, 184 leaves : ill. ; 28 cm.
## Subjects- Subjects / Keywords:
- Alloys ( jstor )
Aluminum ( jstor ) Geometric angles ( jstor ) Geometric series ( jstor ) Material concentration ( jstor ) Nucleation ( jstor ) Particle size distribution ( jstor ) Solid solutions ( jstor ) Wave diffraction ( jstor ) Zinc ( jstor ) Aluminum ( lcsh ) Dissertations, Academic -- Materials Science and Engineering -- UF Materials Science and Engineering thesis Ph. D Precipitation hardening ( lcsh ) X-rays -- Scattering ( lcsh ) - Genre:
- bibliography ( marcgt )
non-fiction ( marcgt )
## Notes- Thesis:
- Thesis--University of Florida.
- Bibliography:
- Bibliography: leaves 177-183.
- Additional Physical Form:
- Also available on World Wide Web
- General Note:
- Typescript.
- General Note:
- Vita.
- Statement of Responsibility:
- by John Thomas Healey.
## Record Information- Source Institution:
- University of Florida
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- University of Florida
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- Copyright [name of dissertation author]. Permission granted to the University of Florida to digitize, archive and distribute this item for non-profit research and educational purposes. Any reuse of this item in excess of fair use or other copyright exemptions requires permission of the copyright holder.
- Resource Identifier:
- 000175126 ( alephbibnum )
03020498 ( oclc ) AAU1595 ( notis )
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GUINIER-PRESTON ZONE EVOLUTION IN 7075 ALUMINUM by JOHN THOMAS HEALEY A DISSERTATION PRESENTED TO THE GRADUATE COUNCIL OF THE UNIVERSITY OF FLORIDA IN PARTIAL FULFILLMENT OF THE REQUIREMENTS FOR THE DEGREE OF DOCTOR OF PHILOSOPHY UNIVERSITY OF FLORIDA 1976 UNIVERSITY OF FLORIDA 3 1111111111111111111111111111111111111111111262 08552 326111111111111111111 3 1262 08552 3263 To Connie ACKNOWLEDGMENTS The author wishes to thank his chairman, Professor Robert W. Gould, who devoted much of his time, and gave both many helpful suggestions and encouragement during the course of this study. He is also indebted to other mem- bers of his committee, Professors John J. Hren, Robert E. Reed-Hill and Frank Blanchard, as well as Professor Rolf E. Hummel, whose valuable assistance and suggestions have aided in the completion of this work. The author is grateful to the National Science Foundation for the financial aid of NSF grant # GH 31817. The author also wishes to thank John M. Horn, William Schulz, John Garrett and Vicki Turner for their assistance provided during this endeavour. iii TABLE OF CONTENTS ACKNOWLEDGMENTS............... LIST OF TABLES............. LIST OF FIGURES............. ABSTRACT................... INTRODUCTION ................ CHAPTER REVIEW OF PREVIOUS WORK................ THEORY OF SMALL ANGLE X-RAY SCATTERING. EXPERIMENTAL EQUIPMENT AND PROCEDURES.. The Alloy............................... Small Angle and Tensile Test Sample Preparation............................. Resistometry Sample Preparation......... Solutionizing and Aging............... Wide Angle Diffraction................. Small Angle X-Ray Scattering........... Resistometry Measurements.............. Tensile Testing......................... EXPERIMENTAL RESULTS................... Resistivity............................. Small Angle X-Ray Scattering . .... .. . .. .. ..... 6 7 Preliminary Studies of Initial Aging.........67 I II III IV ............. ............. ............. ............. ............. PAGE ....vi ...vii .... xi .....4 ....23 .... 48 ....48 .... 48 ....49 .... 50 ....51 ....53 .... 62 ....63 ....64 ....64 PAGE Change in Zone Size Parameter....... ........68 Relative Integrated Intensity...............110 Wide Angle Diffraction......................... 110 Metallography..................................131 V DISCUSSION OF RESULTS........................... 132 Small X-Ray Scattering.......................... 132 CONCLUSIONS ............................................152 APPENDIX ............................................ ... 155 BIOGRAPHICAL SKETCH....................................177 BIBLIOGRAPHY ...........................................178 LIST OF TABLES Table Page 1 Equilibrium Phases.............................. 20 2 Machine Parameters .............................57 3 Summary of Aging Series........................ 67 4 Summary of Results of Wide Angle Diffraction..120 5 List of Variables Used in SAXS Program........157 6 Data Input for SAXS Program................... 160 LIST OF FIGURES Figure Page 1 Reciprocal Space of the Aluminum Matrix and Guinier-Preston Zones ......................26 2 Direction of Incident and Scattered Beam........32 3 Moment Distribution of a Particle About the Origin .....................................33 4 Geometric Relation of a Particle to the Origin .........................................36 5 Schematic Representation of Guinier De Wolff Camera...................................52 6 Schematic of Kratky Camera from Interna- tional Union of Crystallography Commission on Crystallographic Apparatus.................. 54 7 Schematic of Slits System in Kratky Block......55 8a Scattering Curve for Lockheed Pyrolitic Graphite .......................................59 8b Porod Region of Lockheed Pyrolitic Graph- ite Scatterer...................................61 9 Resistivity Aging Behavior of 7075 Wires.......65 10 Arrhenius Plot of Resistivity Maxima............66 11 Fixed Angle Scattered Intensity for 7075 A lloy ..........................................69 12 Arrhenius Plot of Time to Reach Scattered Intensity of 1.5 Counts Per Second.............70 13 Preliminary Study of Effect of Deforma- tion-Fixed Angle Scattered Intensity........... 71 vii Figure Page 14 Typical Small Angle Scattering Curve of Directly Aged Sample ........................... 73 15 Typical Small Angle Scattering Curve from a Preaged Sample............................... 75 16 Evolution of Guinier Radius as Determined from SAXS for A, B, D Series................... 77 17 Evolution of Guinier Radius as Determined from SAXS for E, F, FlH Series................. 78 18 Evolution of Guinier Radius as Determined by SAXS for G and H Series..................... 80 19 Evolution of the Guinier Radius as Deter- mined by SAXS for J and K Series............... 82 20 Evolution of Guinier Radius as Determined by SAXS for L Series........................... 84 21 Evolution of Guinier Radius as Determined by SAXS for M Series........................... 85 22 Porod Radius for A, B, D Series as Deter- mined by SAXS ..................................88 23 Porod Radius Evolution for E, F, Flh Series as Determined by SAXS............................90 24 Porod Radius Evolution for J and K Series as Determined by SAXS............................91 25 Porod Radius Evolution for L Series as Determined by SAXS...............................93 26 Porod Radius Evolution for M Series as Determined by SAXS...............................94 27 Radius of Maximum Abundance for A, B, D Series as Determined by SAXS...................95 28 Radius of Maximum Abundance for E, F Series as Determined by SAXS............................96 viii Figure Page 29 Radius of Maximum Abundance for F1H Series as Determined by SAXS............................97 30 Radius of Maximum Abundance for J and K Series as Determined by SAXS...................98 31 Radius of Maximum Abundance for L Series as Determined by SAXS............................99 32 Radius of Maximum Abundance for M Series as Determined by SAXS ......................... 100 33 Particle Size Distribution Evolution for A Series ...................................... 101 34 Particle Size Distribution Evolution for B Series ......................................102 35 Particle Size Distribution Evolution for D Series ......................................10 3 36 Particle Size Distribution Evolution for E Series ......................................104 37 Particle Size Distribution Evolution for F1H Series ....................................105 38 Particle Size Distribution Evolution for J Series ......................................106 39 Particle Size Distribution Evolution for K Series ......................................107 40 Particle Size Distribution Evolution for L Series ......................................108 41 Particle Size Distribution for M Series.......109 42 Relative Integrated Intensity for A and B Series ......................................111 43 Relative Integrated Intensity for E, F, F1H Series ....................................112 Figure Page 44 Relative Integrated Intensity for G and H Series ......................................114 45 Relative Integrated Intensity for J and K Series ......................................115 46 Relative Integrated Intensity of L Series.....117 47 Relative Integrated Intensity for M Series ....118 48 Densitometer Scans of Eseries (1250C Age) Guinier-De Wolff Films........................ 119 49 Results of Tensile Tests for A, B, D Series...122 50 Results of Tensile Tests for E, F Series ...... 123 51 Results of Tensile Tests for G, H Series...... 124 52 Results of Tensile Tests for J Series.........126 53 Results of Tensile Tests for K Series ......... 128 54 Results of Tensile Tests for L Series.........129 55 Results of Tensile Tests for M Series.........130 56 Schematic Representation of Type B Ser- rated Flow.....................................143 57 Interaction of a Dislocation Line With the Coherency Strains Surrounding a Guinier- Preston Zone...................................145 58 Representation of Area Created by Passing of a Dislocation Through a Guinier-Preston Zone, Along with the Applicable Geometry...... 147 Abstract of Dissertation Presented to the Graduate Council of the University of Florida in Partial Fulfillment of the Requirements for the Degree of Doctor of Philosophy GUINIER-PRESTON ZONE EVOLUTION IN 7075 ALUMINUM by John Thomas Healey June, 1976 Chairman: Dr. Robert W. Gould Major Department: Materials Science and Engineering This research has been concerned with the growth of Guinier-Preston zones, their transformations, and the re- lation of these substructures to the mechanical properties of 7075 aluminum. Thermomechanical treatments, as well as thermal preage treatments, were performed on the alloy. The principle technique of this investigation was X-ray small angle scattering. Mechanical treatments did not change the aging se- quence, but did increase the aging kinetics. Preage treatments produce a narrow stable Guinier-Preston zone size distribution which remains stable upon final ele- vated aging. Wide angle scattering indicated that no large misfit existed between the matrix and the Guinier-Preston zones, and that the zones transformed to an intermediate precipi- tate n', than to the equilibrium phase n. xii INTRODUCTION The phenomenon of age hardening in aluminum alloys was first observed by Wilm in 1905. The discovery was quite accidental. Wilm had been trying to harden an alu- minun-4% copper-0.5% silicon-0.5% manganese alloy by quenching from 525C in a manner analogous to steels. The experiment initially was a failure as the material proved to be quite soft, but several days later, while measuring the hardness, Wilm found it had hardened at room temper- ature. Wilm could not explain this increase, as optically he could not observe any structural differences between the hard and soft specimens. Wilm did not publish this work until 1911, at which time he merely presented the results without attempting to offer an explanation. The first explanation of this anomaly was offered by Mercia, Waltenberg and Scott, who concluded that since the solid solubility of copper and silicon is higher at 5250C, the increase in strength must be due to the slow precipi- tation of some second phase at room temperature. Jeffries and Archer further proposed that these second phases were really submicroscopic particles of the equi- librium phase. The proof of the nature of the hardening agents-had to wait until 1938, when Guinier and Preston first demonstrated that small clusters of solute atoms would produce diffuse x-ray scattering at small angles. This technique provided the means to directly study these coherent clusters, today known as Guinier-Preston zones or merely G.P. zones. The next major advance in the study of precipitation was the continuing development of the transmission elec- tron microscope and the development of sample preparation techniques. This enables the researcher to actually see the zones, their distribution within the matrix, and to follow their evolution with time. Transmission electron microscopy also offered the ability to observe the rela- tionship of the zones to defects in the matrix, such as dislocations and grain boundaries. Extensive studies into many of the age hardenable systems have been made, employing a wide variety of tech- niques; hardness measurements, resistometry, small angle x-ray scattering and, somewhat later, transmission elec- tron microscopy. The basic goal of all these investi- gations has been to obtain an insight into the complex mechanisms of age hardening, and, thereby, be able to predict alloy mechanical properties by a detailed knowl- edge of the zonal state. Polmear and Gerold have done much of the early work on the aluminum-zinc-magnesium system. These researchers have studied the effects of various alloy compositions, and, in conjunction with these, many different aging se- quences. The next logical step in the investigation of the age hardening process is to study the complex ternary and quaternary commercial alloys. This is the subject of this dissertation. The present investigation is con- cerned with characterizing the various zonal states of the commercial 7075 alloy (5.7% Zn-2.6% Mg-1.6% Cu), and their interrelation with mechanical properties. The main technique of the investigation will be x-ray small angle scattering. CHAPTER I REVIEW OF PREVIOUS WORK Extensive work has been done on the decomposition of aluminum-zinc alloy systems. Some of the earliest work on this system was done by Guinier, who showed that the clusters which formed homogeneously upon quenching were spherical in shape. The rapid growth of the clusters, or Guinier-Preston zones, has been attributed to enhanced diffusion in a system supersaturated with quenched-in va- cancies. Resistometry work of Panseri and Federighi2 has shown the formation energy for vacancies to be 0.70 eV, the activation diffusion energy to be 0.54 eV, and the binding energy between zinc atoms and vacancies to be on the order of 0.05 eV. Using small angle x-ray scattering, Gerold3 has shown that the spherical zones continue to grow and main- tain a spherical shape until a Guinier Radius of approxi- mately 35 A is reached. At this size, growth is arrested and the zones transform to an elliptical shape before growth recommences. This shape change is a result of in- ternal rhombohedral straining whose anisotropy causes the elliptical shape. 4 Gerold has further shown that the composition of the zones can be determined from the integrated intensity. In this manner, the limits of the metastable miscibility gap have been calculated. Carpenter and Garwood have shown the aging se- quence in Al-Zn alloys to be: Spherical G.P. zones Elliptical G.P. zones - Rhombohedral a' Cubic a' e (hexagonal zinc) The rhombohedral a' phase is a transition phase resulting from a partial loss of coherency of the zones during growth. The rhombohedral a' phase retains coherency with the (111) plane of the matrix. Further loss of coherency results in the formation of the cubic transition a' phase. Both of the a' transition phases remain homogeneous throughout the matrix as they are both nucleated from the homogeneous Guinier-Preston zones. The final equilibrium phase, e (hexagonal zinc), is nucleated heterogeneously on dislocations only during high temperature aging or pro- longed aging of the a' phases. The commerical usefulness of the aluminum-zinc binary alloys is practically non-existant as these alloys will overage at temperatures below room temperature, thereby eliminating the strengthening effect of the Guinier- Preston zones. The addition of magnesium to the binary aluminum-zinc system greatly alters the aging process of the system. The most important change in the aging character is the ability to retain the enhanced mechanical properties at room temperature. This feature, combined with the fact that the aluminum-zinc-magnesium system provides the highest strengths attainable of age hardenable aluminum alloys, has lead to a considerable amount of research into their aging behavior. Polmear6-11 did much of the early work on the Al-Zn- Mg system. Polmear used hardness measurements to follow the aging sequence. By plotting the incubation time of the hardness curves, that time during which no increase in hardness is observed, versus the reciprocal of the aging temperature, Polmear produced a series of c curves whose intersections indicated the equal probability of the existence of two phases. Using this technique, Polmear constructed the surface of the upper temperature limit of stability for Guinier-Preston zones.11 For an alloy of 5.7 / Zn-2.6 / Mg, he shows this upper temperature to be approximately 1750C.1l Polmear explains the greater stability of Guinier-Preston zones with increasing Mg con- tent as the greater ability of Mg + Zn to retain coheren- cy with the aluminum lattice.7 The explanation for the clustering of Mg and Zn together can be explained thermo- dynamically. A mixture of aluminum and zinc, or of alu- minum and magnesium, show a positive heat of solution, while a mixture of magnesium and zinc possesses a negative heat of solution. For this reason, the magnesium and zinc atoms will be expected to cluster together on quenching. A minimum in the heat of solution should be obtained at some fixed Mg:Zn ratio. Polmear1 states that the con- tinuous nature of the Guinier-Preston zone surface indi- cates that the composition and structure of the zones must remain fairly constant over a wide variation in composi- 12-13 tion. Gerold, however, has shown that the composi- tion of the zone is dependent on the Mg:Zn ratio. If the ratio is greater than unity, the zones have a MgZn compo- sition, and if the ratio is less than unity, the composi- tion becomes Mg3Zn5. The continuity of Polmear surface can be explained by the fact that there is not an abrupt change in zinc composition at these particular ratios, but rather a gradual transition between the two compositions, and that the thermal response of the two are very similar. Nicholson and Lorimerl4 used transmission electron microscopy to determine the upper limit of Guinier-Preston zone stability. They quenched their samples directly from the solutionizing temperature to the aging temperature, and observed the distribution of phases present. Since Polmear had already shown that nucleation of Guinier- Preston zones was homogeneous while nucleation of n from the solid solution was not, Nicholson and Lorimer4 des- cribed the temperature at which the nucleation switched from homogeneous to heterogeneous as the upper temperature limit of Guinier-Preston zones. The limit determined in this manner was approximately 150C lower for a 5.9 / Zn- 2.9 W/ Mg alloy.14 Lorimer's procedure yields the upper temperature limit for nucleation of zones, while Polmear's gives the upper limit for the existence of zones nucleated at a lower temperature. Mondolfo, Gjosten and Levinson proposed the se- quence of aging in the Al-Zn-Mg system to be: spherical G.P. zones -+ -+ Embury and Nicholson,16 as well as Tomita et al.,17 have shown that except for very low concentration (>.5 atomic % Mg), the G.P. zones remain spherical throughout the aging 17 process. Tomita observed the diffuse scatter at the base of the (111) peak in 6.8 atomic % Zn and .06 2.0 atomic % Mg alloys, and concluded that when this size effect scatter disappeared, the strain in the matrix caused by the mismatch between the zones and matrix had become negligible. Without the rhombohedral strains pre- sent as in binary Al-Zn alloys, the zones will not undergo a shape change. Graf18 has shown that the transition phase n' grows on the (111) planes of the aluminum matrix, although some controversy exists as to the mode of nucleation of the phase. Asano and Hirano9-21 propose from their calo- metric studies, that Guinier-Preston zones are an inde- pendent aging product, and do not act as a nucleation 22 8,10-11 site for the formation of r'. Parker,22, Polmear,' Thompson23 and Baba24 indicate that the presence of addi- tional elements, particularly silver, chromium and copper, aid in the heterogeneous nucleation of n'. Dinkeloh, Kralik and Gerold,25 as well as Lorimer and Nicholson,14 support the theory that Guinier-Preston zones larger than a certain critical size can transform directly to the transition phase; the support for this being the homo- geneous distribution of q'. The work of Chang and Morral2 indicates that while the Guinier-Preston zones do not transform directly to n', zones of sufficient size can act as nucleation sites for the heterogeneous nucleation of the precipitate. The homogeneous distribution of n' often observed would be explained by this process, as the distribution of critically sized zones which could act as nucleation sites would be homogeneously distributed through the matrix. Bardhan and Starke27 give the criti- cal Guinier Radius necessary for Guinier-Preston zones to act as nucleation sites at 35 A, which is curiously simi- lar to the size at which zones change from spherical to ellipsoidal in the binary Al-Zn alloys. While the n' phase is generally accepted to be hexagonal and semico- herent with the aluminum matrix,15,26 with lattice para- 0 0 77 28 77 Ryum28 finds n' to be mono- meters a =4.96 A, co=8.68 A, Ryum finds q' to be mono- clinic with lattice parameters a=b=4.97 X, c=5.54 X and 29 8=120. Thackery,2 on the other hand, finds no evidence for the existence of the n' phase, and proposes the aging process to be: SOLID SOLUTION G.P. zones n while speculating that previous x-ray investigations may have mistaken his X phase for n'. The equilibrium n phase (MgZn2) can be formed dir- ectly by the change of lattice parameter and.loss of co- herency of '.15 The MgZn2 phase can also be nucleated heterogeneously at dislocations and grain boundaries.26'29 With Mg:Zn ratios greater than 3:7,22 the n phase will it- self transform to the equilibrium T ((Al,Zn)49Mg32) phase with aging treatments above 200C; although, it has been shown by Schmalzried and Gerold12 that this transformation shown by Schmalzried and Gerold that this transformation remains incomplete, even with extended aging at tempera- tures greater than 200C and is only completed with aging above 3000C. Ryum2830-31 shows that this aging sequence is only one of a possible three which can occur simultaneously or independently, depending upon the conditions. The possi- ble aging sequences as seen by Ryum are: SOLID SOLUTION n SOLID SOLUTION "nuclei" n' n SOLID SOLUTION G.P. zones n' n The new "phase" introduced by Ryum30 ("nuclei") is defined as a vacancy-solute atom aggregate, but with a higher upper temperature limit. Ryum31 presents the upper tem- perature limit of the nuclei to be 300C as opposed to 1800C for Guinier-Preston zones. The first of Ryum's se- quences (solid solution -* ) would have a high activation 31 energy and would occur heterogeneously on dislocations and grain boundaries, or homogeneously, with a low parti- cle density, after long aging periods. Bardhan and 29 Starke also show the presence of quench-in MgZn2 in a 5.8 weight % Zn-.71 weight % Mg, but not in a .51 weight % alloy. They attribute this difference to a greater degree of quench clustering in higher magnesium alloys. 31 Ryum31 states that the second sequence (solid solu- tion + nuclei -- nr) occurs with quenching to room temperature, holding for five seconds, then aging at 1500C. The nuclei were not observed in the transmission electron microscope, with n' lying on {111} matrix planes being the first observable phase. This reaction did not occur in the grain boundary region or adjacent to dislo- cations in the matrix due to the low vacancy concentra- tion. The final sequence is identical to that presented previously. Several authors12'16'2931 have presented the orien- tation relationship of the MgZn2 phase andthe transition phase q' with the matrix. Mondolfo et al.15 suggest that the basal plane of the hexagonal q' phase lies on the (111) plane of the aluminum with the following relation- ship: (0001) n',(111)Al, (1120)1 (110)A1 12 Schmalzried and Gerold 12 have proposed three possible orientation relationships between n phase and the aluminum matrix: a.) (1010) nI(100)Al, (0001) 1 I(011)A b.) (1010) n (110)A1, (0001)n (111)A1 c.) (1010)n| (121)A1, (0001) I(111)Al Thackery29 has shown that Gerold's relationship c.) can be written as: (1120) n (i11(0)A, (0001)n | (Il)A which would make this orientation identical with that dis- covered by Mondolfo for n', thereby indicating a direct transition between the two phases. Embury and Nicholsonl6 give the orientation relationship to be: (0001) n|(110)Al, (1010) nl (110)A which presents a larger mismatch between the two phases 12 29 than those presented by Gerold.2 Thackery,29 examining an aluminum-6 weight % Zn-2 weight % Mg alloy, determined six possible relationships of n to the aluminum matrix. 1.) (1210) II(111)A1, (0001)n (110)A 2.) (1210) n (111)A1, (3032) In(110)A 3.) (1210) 1(111)Al, (2021) 1(121)Al 4.) (1210) T (111)A1, (1014) (110)A1 5.) (0001)n I(111)Al, (1010) |(110)A1 6.) (1010) |I(110)A, (0001)n (001)A1 Type 1 orientations were observed by Thackery to be laths lying on {111} planes of the matrix with Types 2, 3 and 4 basically similar with a rotation of the precipitate on the (111) planes. Type 5 corresponds to Gerold's Type b.), while Type 6, which Thackery distinguished as elon- gated eight-sided platelets, is identical to Gerold's Type c.). Thackery found no evidence for the existence of Gerold's Type a.), or of the Embury-Nicholson orientation relationship. Thackery's Type 1 precipitate is presented as the most common, being nucleated heterogeneously on dislocations and atgrain boundaries and forming as plate- lets on (111) matrix planes. Type 5 Thackery represents as forming only as coarse particles resulting from direct high temperature quenches, while Type 6 constitutes the majority of particles in samples first quenched to room temperature then aged.29 28 Ryum found evidence for the existence of all three orientations proposed by Gerold, Embury and Nicholson's 28 relationship, and those of Thackery. Ryum28 also noted that Gerold's orientation Type a.) had an appearance simi- lar to that of Thackery's X phase.29 Resistometry has been widely used to calculate the formation and activation energies of vacancies and vacan- cy complexes.2'14'23'32-39 Panseri and Federighi2'34 did much of the early work with resistivity on the aluminum- zinc-magnesium system. By measuring the time to reach a maximum of resistivity as a function of aging temperature, they were able to calculate the values for the binding en- ergy of zinc vacancy couples to be .06 eV and that of mag- nesium vacancy couples to be .54 .08 eV, a value also given by Gould and Gerold.40 They also concluded that magnesium vacancy couples can diffuse freely at room tem- perature, and that the movement of zinc atoms occurs by means of these couples, and not via single vacancies as in binary aluminum-zinc alloys. 33,41 Perry, by employing the initial rate of aging as well as the time to peak resistivity, has determined the apparent vacancy formation energy to be .70-.71 eV. In working with concentrated (i.e., 10 weight %) zinc alloys, Perry41 contends that the zinc vacancy binding energy measured is not the true binding energy, but merely an effective value obtained due to the clustering of zinc atoms on quenching; the value Perry obtained was .06 eV, in agreement with Panseri. Perry, however, obtains a vastly different binding energy (.17 eV) for the magnesium vacancy complexes, using a .009 atomic % Mg-4.01 atomic % Zn alloy.35 Perry also notes that this small addition of Mg was sufficient to slow down the aging process due to the trapping of the vacancies by magnesium atoms. Panseri and Federighi, as well as Gould and Gerold, note that magnesium initially slows the aging process by trapping vacancies, but also extends the aging time by up to an order of magnitude. This effect is due to the fact that magnesium will act as a vacancy source when the matrix has become depleted in vacancies. 37 Ceresara and Fiorini,37 using a similar resistometric technique, have measured the total kinetic activation energy in Al-5 weight % Zn-l and 3.4 weight % Mg alloys to be .67 eV, independent of magnesium concentration. Several authors9-21'4243 have made use of a calori- metric method to study the thermal stability of Guinier- Preston zones, their reversion, and transformations to 19-21 transition and equilibrium phases. Asano and Hirano have shown that Guinier-Preston zones become stable to higher temperatures with longer aging times at lower tem- peratures. Their calometric scans indicate that with short aging times, the reversion of the zones is complete before formation of n', which would indicate heterogeneous nucleation of the transition phase, independent of the zones. However, with longer low temperature aging times, the heat absorption of zone reversion overlaps the heat evolution of n' formation, which would imply some sort of n' nucleation on the zone sites. Of particular importance to commercial alloys is their quench rate sensitivity, especially due to the pre- sence of secondary alloying elements. Several workers22-2344-46 have studied the quench rate sensitivity of Al-Zn-Mg alloys with the additions of chro- mium, silver, zirconium, manganese and copper. The con- sensus of these authors is that all these secondary alloy elements produce quench rate sensitivity, roughly in the order presented. The cause, however, is not the same for all these elements. In the case of copper, which presents the least effect, Mondolfo47 has shown that the copper precipitate (CuMgAl) can precipitate with a lower activa- tion energy and, being structurally similar to MgZn2, can act as easy nucleation sites for this phase; much the same 23 is true for zirconium (Al2Zr),23 and also for silver which 42 clusters readily in aluminum alloys.42 The effects of chromium and manganese are more complicated. Both of these elements have a low solubility in aluminum and would 46 tend to form clusters, but more important is that both elements act to retard recrystallization by forming sub- grain structure.46 This structure acts as both vacancy sinks and as additional heterogeneous nucleation sites, the former removing vacancies from the aging sequence and the latter, solute atoms. Both of these effects will be detrimental LO the final mechanical properties. The most important parameter affecting the mechanical properties is the size and distribution of the prepre- cipitates and the precipitates. The only two ways to mea- sure these parameters directly are small angle scattering and transmission electron microscopy. Many investigators have used TEM measurements to determine the precipitates 46 present,46 their orientation relationship with the matrix12'1427'293031 and distribution within the matrix.26'4853 The lack of resolution during the early stages of aging and the tediousness of measurement of G.P. zone size using transmission electron microscopy makes x- ray small angle scattering the most accurate and fastest method of direct measurement of the precipitation process. Guinier5456 and Gerold5759 developed much of the theory as it relates to the solid solution; this will be presented later. In addition, Gerold12-13'25'58-60 also did much of the early work on the Al-Zn-Mg system. Tomita,17 as well as Bardhan and Starke,27 have shown that except for the very low concentration alloys, there is no zone size arrest, as in the Al-Zn system. Harkness, Gould and Hren4850 have developed the mathematics for converting the measured experimental small angle scat- tering parameters into actual particle size distributions; their development will be presented elsewhere. Most of the direct measurements of G.P. zone size distributions as a function of aging conditions have been done on pure research grade alloys. The next logical step would be to measure their evolution in a commercial alloy complicated by the presence of a large number of alloying agents and also by the presence of impurity atoms. The purpose of this work will be to measure quantitatively the evolution of Guinier-Preston zones in a commerical 7075 alloy, the effects of thermomechanical treatments on the zones, and their relationship to the mechanical proper- ties. The main technique of this investigation will be x-ray small angle scattering. Equilibrium Phases The equilibrium phases possible in the aluminum-rich corner of the aluminum-zinc-magnesium system are given in Table 1. The phases possible involving the trace impurity and alloying elements present, such as Si, Mg, Ti, Cr and Mn, will not be considered as their concentrations are suffi- ciently low that they will be either incorporated in solid solution, miscible in one of the other equilibrium phases, or so widely dispersed that they will have little effect on the alloy as a whole. Table 1 Equilibrium Phases Phase Notation Crystal Lattice Parameter A Structure Al a FCC a =4.049 (47) MgZn2 n HCP ao=5.16-5.22 (47) CuMgAl c0=8.49-8.55 (47) (Al,Zn)49Mg32 T cubic ao=14.29-14.71 (47) CuMg4Al6 ?X HCP a =2.66 (29) 0 c =4.94 o FeAl3 X Monoclinic a=15.48 b=8.083 c=12.17 B=107043' The a phase lattice parameter varies linearly from a low of 4.018 A at 57 weight % zinc to a maximum of 4.1294 at 17 weight % magnesium. The n phase, MgZn2, is a proto- type of the hexagonal Laves Phase.47 Thomas and Nutting88 have shown that an appreciable amount of copper can also be dissolved in the n phase. Mondolfo. Gjosten and Levinson5 have shown that the r phase is the only ternary phase that occurs with aging temperatures below 2000C. The T phase, (Al,Zn)49Mg32, is a complex cubic struc- 47 ture of 162 atoms to the unit cell. The lattice parameter varies as the Al:Zn ratio with the magnesium content remaining essentially constant. Little, Hume-Rothery and Raynor89 have shown that the T phase can dissolve a significant amount of copper also, as the CuMg4Al6 phase is completely miscible with T. Mondolfo et al.1 have shown that the T phase is slow to nucleate and is present only with extended aging above 2000C. Thackery29 indicates an undetermined phase present in Al-6 weight % Zn-2 weight % Mg, and designates it as X phase. He determined it to be hexagonal, appearing as triangular-shaped particles whose sides are parallel to (110) matrix planes, and whose lattice parameters are ex- tremely similar to those of pure zinc. The X phase may truly be zinc precipitated from small zinc clusters with limited quantities of alloying elements in solid solution. The X phase is the result of iron being an impurity in all commercial aluminum alloys. It is very slow to homogenize and does not tie up any of the major alloying elements, and is subsequently important to the aging pro- cess only to the degree that it can act as a heterogeneous nucleation site. The real importance of the X phase lies in its detrimental effect on the fracture mechanics of the alloy. 22 The chromium addition is present to enhance stress corrosion resistance and the phases formed do not greatly effect the aging process. Titanium is added as a grain boundary refiner, and the other impurities present may affect the kinetics of the aging process, but do not greatly tie up the effective alloying constituents. CHAPTER II THEORY OF SMALL ANGLE X-RAY SCATTERING Small angle scattering has a variety of applications; the study of macromolecules, polymers and clays. In 1938, Guinier proposed their use for the study of inhomogeneities in dilute alloy systems. There are sev- eral restrictions placed on the study of dilute alloy systems. First of all, the alloy must contain regions of electron density inhomogeneities of a size greater than atomic radii. Secondly, these inhomogeneities must be smaller than 1000 X, the approximate upper limit of res- olution at that time. The inhomogeneity must have a fairly significant electron density difference from the matrix and must be reasonably concentrated (approximately 1% of scatterers for Guinier-Preston zones) in order to produce measurable scatter in a thin sample necessitated by x-ray absorption in the sample. (The thickness for an aluminum alloy using MoKa radiation is approximately 0.7 mm.) The last of these restrictions, a measurable scat- ter, is largely restricted by instrumental parameters. As the electronics become quieter and more sensitive, the minimum quantity of inhomogeneities also becomes lower. This is a result of the statistics of the count rate. To understand the origin of the scatter, the concept of the reciprocal lattice must first be understood. In reciprocal space, all lattice planes (hkt) of real space are represented as reciprocal lattice nodes. The size, shape and spacing of the reciprocal lattice nodes are inversely proportional to the size and shape of the crys- tal in real space, and the interplanar spacings, res- pectively. Thus, large crystals in real space would be represented as small point-like nodes in reciprocal space, while large platelets would appear as long narrow cylin- ders in reciprocal space. In the case of Guinier-Preston zones, there are two sets of reciprocal lattices; the matrix and the G.P. zones themselves. The matrix, being large crystals, will be re- presented by sharp reciprocal lattice points, neglecting stress and temperature effects. The G.P. zones will have their reciprocal lattice nodes centered at the same posi- tion as the matrix nodes, as the G.P. zones have the same crystal structure as the matrix. However, the small size of the G.P. zones produces reciprocal lattice nodes which are broad and spread. Applying the Ewald Sphere concept, diffraction will occur whenever the diffraction beam vector, (IS =~), intersects a reciprocal lattice node. The diffraction vector, S, is drawn from the center of the Ewald Sphere, S whose radius, |-1I, is directed to the (000) node, to any reciprocal lattice node that intersects the Ewald Sphere. When this happens, diffraction is said to occur. In the case of large crystals where the nodes are sharp and point-like, diffraction is restricted to these points, with nothing occurring between them. This diffraction is restricted to definite angular regions. However, with small G.P. zones, the nodes are large; thus,-there is a greater probability of the Ewald Sphere intersecting a reciprocal lattice node of the G.P. zones surrounding those of the matrix. Figure 1 shows reciprocal space of the matrix and the G.P. zones. This accounts for the greater angular range of diffraction for the G.P. zones. The two diffraction peaks superimpose and the G.P. zone peak is evidenced as a deviation from a Gaussian distri- bution at the base of the matrix peak. This scatter is present around all the Bragg peaks, as well as the (000) node. Observation of this scatter is generally simplified around the (000) node, as such effects as strain and tem- (1) C-,, L&j NN ccc 0 E I U) 0 Gd 4 I ca :2: 0 o 0 04 r-4 0 r4 P-4 perature which cause broadening of the Bragg diffraction peaks are not present around the (000) node. The inten- sity is also maximized around the (000) node as the atomic scattering factor is maximized here, and the Compton Scat- tering, which subtracts from the intensity (as (1-cose)), is minimized. In order to understand the small angle scattering that results from the Guinier-Preston zones, it is nec- essary to start with the first principles of x-ray scat- tering. A dilute system is one in which there is a large volume available per particle as compared to the total volume of all the small particles. In addition, thin samples will only be considered such that absorption ef- fects can be neglected. The basic development of 1,2 Guinier'2 will be mostly followed. For the case of a small isolated crystal, such as a G.P. zone, Guinier has shown that the scattering power per unit cell is: F2 Fhk 2 I(s) hk EE E (s-rhk c hh h where: V is the volume of the diffracting crystal V is the volume of a unit cell Sis the structure factor of a particular node Fhkz is the structure factor of a particular node E(s) is the Fourier transform of the form factor rhk, is the structure factor of a particular node 0 S-- Around the (000) node, the vector rhk =0. Thus, around the (000) node: 2 F(000) 2 I(s) = v0 |E(s) | c For very small angles, the structure factor, F(000), is equal to the number of electrons (N) present in the unit cell. The average electron density (p) of the unit cell is then: (000) N V V c c The total scatter (IN(s)) from all electrons in the crys- tal of total volume V is given by Guinier5 as: V 2 2 c where: V- is the total number of unit cells in the crys- c tal Thus, the maximum scattered intensity occurs at s=0, where E(0) = V. 2 2 Ima = P V This maximum intensity peak of large crystals is masked, This maximum intensity peak of large crystals is masked, however, by the main unscattered beam. Since s=2sinO 20 c1 where 0 is small, s=---- Also, s= ; therefore: d The minimum angle where scatter can be observed is e=10-3 radians. This would correspond to a particle diameter (d) 0 of approximately 1000 A. A Guinier-Preston zone can be considered as an iso- lated crystal within a medium other than vacuum, as long as it is a dilute system, as defined previously; the in- tensity from each particle will merely add. If the med- ium electron density is given as po and the isolated par- ticle electron density as p, the scattering power per par- ticle is: I(s) = (p-po)2 |E(s)|2 54 Guinier has determined that in order to obtain the aver- age scattering power per particle if all particles are identical, 1(s) 2 must be averaged over the entire sphere of radius s=|s|. Thus, 2 2 average value of 1|(s)2 I(s) = (p-po) {over the sphere } of radius s=Is| To obtain the average value of E(s), the form function must be integrated over the volume, and for the case of spheres: E(s) = / exp(2fis-x)dV V 54 Guinier has shown that for a spherical-shaped particle: 4 3 E(s) = I a (3(27sa) for a sphere of radius a. Where the function p(2nsa) is given as: (2~s) = 3sin2as 2ias cos2aas Q(2isa) = 3[in -- 3(2-s (27Tas)3 Thus, the scattering power per particle is given by: I(s) = [(p-po) a3]2 (2as) We now wish to approximate the curvature of the in- tensity curve at the center (s=0). The total scattering power IN(s) for N unit cells of a particle of volume V has 54 been shown by Guinier to be: IN(S) = (p-po)2 I(&)12 where: E(s)=V at s=0 Thus: IN(O) = (p-po) V For an isolated particle, the total number of scattering electrons n=pV. For a particle in a matrix, an effective number of scattering electrons can be defined as: n = (p-po)V at s = 0 The total scattering power per particle, n2=(p-po)V, at 2 2 s=0 decreases with s as n (1-Ks ). K, the curvature at the center of the intensity curve, is related to a simple geometric parameter of the particle, and is independent of the shape of the particle. The approximation of this cur- vature at the center of the intensity curve is known as the Guinier approximation. S is the direction of the incident beam on the scat- tering particle, as shown in Figure 2. As can be seen, for very small scattering angles s' is in the direction D, approximately normal to So, where s=(S-S )/X. s is in the plane of the incident and scattered rays, and |s =- for small angles. Take any vector, x, in real space. Then, -* -4 s x = sxD where xD is the projection of the vector x on D. In order to calculate the function E(s), we must evaluate the inte- gral: E(s) = f o(xD) exp(27is xD)dVx V where o(xD) is the cross sectional area of the particle along a plane normal to D at a distance xD from the origin as shown in Figure 3. If the origin of the coordinate c-o Cd 4-J C/. c ac rl 0 ,- C) -c P- C) 0 C) *-H 0Q b e-- Cl system in real space is choose to be the center of grav- ity of the particles, then the sum of the moments of the particles about the origin is zero. Thus, since ExdV = 0 x and dVx = a(xD)dxD fxD o(xD)dxD = 0 Now we can expand the exponential function in the equation 2 of E(s), neglecting terms of order greater than s since s remains very small within the scattering particle. Now: E(s) = f o(xD) exp(27ris xD)dVx = V / o(xD)dxD + 2nisfxD o(xD)dxD 2r2 s 2fx a(xD)dxD The first term of the above equation is the volume of the particle, while the second term is zero, due to the choice of the origin as shown above. In order to evaluate the 2 third term, an additional term, RD, must be defined: 2 1 2 RD V xD P(xD)dxD RD is defined as the average inertial distance along D to the plane TD passing through the origin and perpendicular to D. Thus: 2 2 E(s) = V 2f s VRD which can be rewritten in exponential form as: 2 222 E() = V exp(-22 s RD) Then, the scattering power per particle: I(s) = (P-Po)2 I|(s)j2 becomes I(s) = (p-p)2V2 exp(-4f 2s2RD) RD, as written here, assumes all particles have identical 2 orientations. For a random orientation of particles, RD 2 2 must be replaced by RD, which is the average value of RD for all directions of D. In order to calculate RD, we must remember that D is merely one coordinate axis of an orthogonal coordinate system. The other two axes, U and V, are shown in Fig- ure 4, along with D. The origin is the center of gravity and any point may be represented as xD, xu or xV. The distance r of any point from the origin is given as 2 2 2 2 r = XD + x + Similarly, R, which is the radius of gyration of a par- ticle about its center of gravity, is given by: 36 Q bIJ 0 41) 0~ 413 w co co 4-4 S U 7< 1 00 r- 0 0 4i '4 r4J w 0 w wr bfl 1:. 2 2 2 2 R = R +R +R D u V where RD, Ru and RV are the inertial distances taken along the three coordinate axes D, U and V, respectively. Now, if we consider the rotation of the three coordinate planes about the origin, as would be the case for random parti- 2 2 2 2 cle orientation, R remains constant and RD, R and R, on D' u V, on the average value, are equal, Therefore: 2 2 3R = R This makes the average scattering power per particle S2 22 2 2 2 A4iT s R 2 A nr R I(s) = n exp(-- 7- ) = n exp(- --2 ) 3 3X The radius of gyration of a sphere is given as R = (3 1/2a R a Thus: 222 2 A4TT a c I(s) = n exp(- --4 ) 5X By taking the logarithm of both sides of the equation: 2 4 2 2 2 in I(s) = n n -41 .4343R2 3XZ Guinier54-55 has shown that the radius of the particles can be calculated from the slope of the line obtained by plotting n Is) versus 2. Thus, R, assuming all parti- plotting Zn I(s) versus E Thus, R, assuming all parti- cles are the same size and randomly oriented in a dilute system, is given as: R = .416X/-a A system does not usually contain all identical spherical particles, but a size distribution is usually present. To correct for this, an integration must be made to account for each size particle: 2 6 4 2 2 2 I(s) = (p-po) Kfn(a)a (l 4s a )da O where n(a) is the number of particles of radius a and the a term is the result of V2=Ka This can be written as: I(s) = (p-po)2N KI where N is the total number of particles and , the average value of a, is obtained by: = f n(a)anda O Rewriting the intensity equation in exponential form: 2 8 58-59 Baur and Gerold determined that by plotting n I(s) versus the scattering angle squared (c ), one obtains a value for the Guinier approximation which is defined as: RG 1/2 Porod61 has also developed an approximation for the scat- tering curve, but his deals with the tail region of the curve, rather than the central region. Porod uses the approximation of the form factor: E(s) = f exp(-2ris x)dV x The scattering intensity may be written as 2 (p-po) a2 i 4a I(s) = 4ra- + 4a sin4Tas 8rr s 7rrs 2 + (-- 1-g) cos ias] S TS If there exists a distribution of spheres of radius aK which varies between al and a2, and the number of spheres of radius aK per unit mass is given by gK, then the inten- sity may be written as: 2 (P-Po) 2 1 I(s) [EgK 4raK + ' 3 s + Ecos4iaKs[...] + Esin2raKs[....]] For the case where (2als-2a2s) is much larger than unity, the two sinusoidal summations go to zero as the positive terms cancel the negative terms. Therefore, for a large 1 s, the only remaining term is that containing For S large s, I(s) is written as Iasym, or the asymptotic val- ues of the intensity. 2 (P-Po) 2 1 S [ gK 4aK asym 8f3 -K aKs ra The total surface of the particles per unit mass is given 2 by S = E gK 4ARK. Then, for the point collimation: (P-Po)2 s asym 87T3 s Another important property of the scattering inten- sity is the total integrated intensity, Q. Guinier and Fournet have shown that the total integrated scattering intensity is given by: V a 2 2 Q = -V(p-p dV = 44fs2 I(s)ds O In the case of a metal alloy, the electron densities, p and p must be computed as the weighted average of the elemental concentration in both the particle and the 57-59 matrix. Then, as shown by Gerold:57 S= -[B + M(ZA ZB) a S=V B + M2(ZA ZB) a where: p is the electron density of the particles p is the electron density of the matrix V is the atomic volume a M1 is the solute concentration in the particles M2 is the solute concentration in the matrix ZA is the atomic number of the solute ZB is the average atomic number of the solvent It is necessary to use ZB because in an alloy, sev- eral solvent elements may be present; therefore, the aver- age atomic number must be used to obtain the electron den- sity. Then, the difference in electron density is given by: (P-po) -[M(ZA ZB) M2(ZA ZB) a -[(M1 M2) (ZA ZB)] a 2 1 2 2 (p-po) V [(M1- M2) (ZA- ZB a If p represents the total volume of particles transformed and cpV the total volume of all particles, we can get the average solute concentration, MA, in the bulk alloy as: cM1 + (1 C)M2 = MA and also the integrated intensity as a function of concen- trations: Q = (1 C)(M1 M2) (ZA ZB)2 a The problem remains now of evaluating the integrated in- tensity from the measured scattered intensity. The inten- sity measured by a radiation detector, E(s), is given by: I nfE I(s)de -d E(s) = --2 4r I nf where the constant factor, e can be written as a 4r constant, K. Then: E(s) = KE I(s)de- d I nf where: K = e 4r - I is the Thompson factor (7.9 x 10-26 cm ) 3 n is the number of atoms per cm f is the surface area of the detector r is the sample to detector distance E is the incident beam intensity I(s) is the scattering power d is the thickness of the sample V is the linear absorption coefficient This equation holds true for a beam of x-rays of point cross section. In order to evaluate the intensity measured from a line-shaped beam, Guinier and Fournet55 have derived the following integration: E(s) = KE de-dl (s2 + t1/2dt o t o where t is a reciprocal lattice parameter related to the length of the line beam. The determination of the inte- grated intensity is accomplished by evaluating the inte- gral: 00 2 2 2? 1/2 Q = 4rfs I(s)ds = 4rr/I(s + t2)2dtds O O 55 Guinier and Fournet have shown that this leads to 2y e-pd 2 Q = rErXd 2TTfsE(s)ds 0 0 where 2y is the width of the detector-receiving slit. In order to evaluate the integral: fsE(s)ds O the Porod approximation for the tail of the scattering curve must be adjusted for a line-shaped beam to: E(s) = ks3 Then: S 0 2 fsE(s)ds = f sE(s) + S E(S ) O O All of the above equations assume a dilute system, i.e., no interparticle interference, monoshaped particles and identical composition of all particles. If the system is not sufficiently dilute, the diffraction maxima will not coincide with s=0, but will be centered at some finite s It is still possible to analyse the curvature at the center of the scattering curve, i.e., the Guinier region, and the tail of the curve, the Porod region. Gerold3 has derived an expression for the Porod radius (R ) by assuming a constant volume fraction of the scattering zones. Combining this assumption with the scattering curve, he concluded that the surface area of the zones (S) computed at large angles should be given by: 2 I(~s48 s3 S = 4TNv (P-Po2 V (p-poN Combining this with the total volume of the particles: cV =4 cV = where: V is the total volume of irradiated material c is the volume fraction of the scattering parti- cles 3 cV Gerold3 solved for the expression by use of the equa- s tion for integrated intensity given previously. He showed that: R 3 1 sE(s)ds p U TTI y I(So) S3 58-59 Baur and Gerold58-59 have shown that this reduces to R Recently, Harkness, Gould and Hren8-50 have shown that the particle size distribution can be obtained from the experimentally determined values of the Guinier radius and the Porod radius. Harkness et al.19 determined that a log normal distribution could be used to represent the particle size distribution. A log normal function has the form: F(X) exp[- ( n ) 2 FT( ~ X n a n a where: p is the geometric mean a is the variance Using the general moment equation for a log normal distri- th bution function, the n- moment is given by: 2 n n 2 r = exp[n 9n p + -n 9n o] Applying this relationship to the results of Baur and Gerold,5859 Harkness et al.49 showed that RG = Re [-5-] and R = Solving these equations for p and a yields: RG pn y = nn R 1.71n- G R p RG a n( ) 9n2 o - 3.5 Using these equations and the assumption of a log normal distribution, it is possible to plot the zone size distri- bution from the experimentally determined values for the Guinier radius and the Porod radius. 49 Harkness et al. have also determined that NV, or the number of zones per cm3, can be determined using the zone size distribution and the volume fraction, Vf, of zones obtained from knowledge of the miscibility gap as follows: 4r 71 V f In this manner, it is possible to plot the actual number of particles within each zone size. Changes in the volume fraction can be monitored as a function of changes in the integrated intensity. Gerold3 has shown that the volume fraction of particles at any time, t, can be given as: Qo(t) f(t) = Qo< f(t ) where: f(t) is the volume fraction at time t f(-) is the volume fraction of the fully aged sample as determined from the phase diagram Qo(t) is the integrated intensity at time t Qo(D) is the integrated intensity of the fully aged alloy 47 Using this relationship, the volume fraction of particles can be accurately determined throughout the aging se- quence, and, therefore, accurate particle size distribu- tions can be determined. CHAPTER III EXPERIMENTAL EQUIPMENT AND PROCEDURES The Alloy The 7075 commercial aluminum alloy was supplied as 3/4 inch alclad plate from the Reynolds Aluminum Company. The alloy was analyzed spectrographically by two x-ray methods. First, it was analyzed using an x-ray energy dispersive system at Oak Ridge National Laboratories, and secondly, it was analyzed using a Norelco* crystal disper- sive x-ray spectrometer. The averaged results of these tests showed the composition to be: Zn-5.80 wt % Cr-.22 wt % Mg-2.62 wt % Si-.19 wt % Cu-1.74 wt % Ti-.04 wt % Fe-0.27 wt % Al-balance Small Angle Scattering and Tensile Test Sample Preparation The plate was prepared for rolling by milling 1/16 inch off of each face in order to remove both the alclad layer, as well as any effects of it in the 7075 substrate. The plate was then cold-rolled on a Fennt rolling mill, *Philips Electronic Instruments, Mount Vernon, New York tFenn Manufacturing Company, Newington, Connecticut with intermediate anneals of 30 minutes at 460C to thick- nesses of .85 mm, .80 mm and .72 mm. These thicknesses produced a constant .72 mm thickness for all samples when the final mechanical treatments of 15%, 10% and 0% were later applied. This large reduction with intermediate anneals was sufficient to break up any large precipitates present in the as-received condition. Specimens for the tensile tests were prepared in the same manner, with the gage section being cut on a Tensile- Kut* high speed milling machine, according to ASTM Speci- fication A370 (substandard size) subsequent to aging. Resistometry Sample Preparation Resistometry samples were prepared by first turning 1/2 inch rods from the 3/4 inch plate. This rod was swaged on a Fenn swaging machine, with intermediate anneals at 4600C for 30 minutes, down to .106 inches dia- meter. This thin rod was then drawn using a wire drawer down to .02 inches. Intermediate anneals at 4600C for 30 minutes were necessary after each draw; however, annealing left the wire too soft to draw. Thus, it was necessary to slightly age-harden the alloy by heating for 10 minutes at 1350C after annealing in order to develop sufficient *Sieburg Industries, Danbury, Connecticut strength. The wire was then made into a coil and voltage leads were spot welded onto it. Solution Heat Treating and Aging The small angle scattering and tensile specimens were solution heat treated at 4600C30C for 2 hours in a verti- cal tube furnace equipped with a large Inconel block for temperature stability. Quenching was accomplished by drop-quenching the specimens and holding the rod into an ice water bath (1/20C). The solutionized samples were stored in liquid nitrogen prior to elevated temperature aging. The resistometry specimens were solutionized in a forced air furnace for 2 hours at 4600C and were quenched by dropping into an ice water bath. The wires were mounted onto a phenolic board immediately after quenching and were then promptly aged. All elevated temperature aging was conducted in a Lauda* constant temperature oil bath, capable of main- taining a given temperature within .10C, containing 50 cs Dow Corning' silicon oil. After aging, and prior to testing, the samples were stored in a refrigerator (o4C). *Brinkman Instruments, Westbury, New York tDow Corning Corporation, Midland, Michigan Retesting of specimens, even after long periods of storage in the refrigerator, indicated that no measurable change in the zone state had occurred, provided some elevated temperature aging had been performed. Wide Angle Diffraction Small strips, approximately 1/8 inch x 5/8 inch x .72 mm, were cut from the small angle scattering samples to be used as samples for wide angle diffraction. The wide angle diffraction was done using an Enraf-Nonius,* Guinier-De Wolff camera. The essentials of this camera are shown in Figure 5. This camera is equipped with an incident beam, elastically curved, quartz monochromating crystal, which focuses the primary and diffracted beams onto the diffraction circle (i.e., the film). All expo- sures were run at room temperature with the chamber evac- uated to eliminate air scatter. The use of monochromatic radiation and an evacuated path eliminates fogging of the film and allows long time exposures to be run. The lower limit of detection with this method is on the order of .05 to .10% of a second phase. This technique was used to determine any quenched-in precipitates and also to follow the sequence of precipitation. Four samples can be run *Enraf-Nonius, Inc., Garden City Park, New York L-2 tj Q simultaneously in a Guinier-De Wolff camera. All expo- sures were run for 48 hours, with machine settings of 45 KV-20 ma. The films were densitometered after devel- opment in order to determine intensities of the weak pre- cipitate diffraction lines. Small Angle X-Ray Scattering The small angle scattering experiments were conducted using a Siemen's* Kratky small angle camera. Figure 6 shows a schematic of the essentials of this type of cam- era. Small angle scattering cannot be observed using standard x-ray diffraction equipment due to the divergence of the primary beam; thus, the need for such a system. Figure 7 is a diagram of the collimation system for the Kratky camera. The first slit, or entrance slit, re- stricts the height of the primary beam. The next set of slits, in the Kratky block, restrict the divergence of the primary beam to include only parallel, or nearly so, rays in the incident beam. These two slit systems greatly re- duce the intensity of the primary beam by severely lim- iting the solid angle of the x-ray beam front observed, but, when aligned, are capable of producing a flat, parallel beam of uniform intensity. The condition of uniform in- *Eastern Scientific Sales Co., Marlton, New Jersey LLU h/ E I K) - rrj E _____. ^' \' N 4I: .I' I3 : -11 ~L < ) 0 o U1 0U 0 p * r- 40 HI-- H U C, CO N Q) 0 4 0 r4 Uo) eCo 0 Sr4 COO0 Ei 0 -I4 r-o> * O KRATKY BLOCK Figure 7. Schematic of Slits System in Kratky Block tensity and "infinite" slit width61 is necessary in order to satisfy the conditions of the Porod equation. For the Kratky camera, the condition of infinite width is ex- pressed as: Z < 2m + a where: m is the height of observation above main beam a is the width of the primary beam at the exit slit k is the length of the primary beam that is of uniform intensity A uniform intensity beam of 3.5 cm was obtained on align- ment of the Kratky camera, while the rear slits were set at 1 cm. A graphite crystal diffracted beam monochromator was used in order to measure only monochromatic radiation, and pulse height selection was employed to eliminate any har- monic diffraction which might occur. Table 2 gives the operating machine parameters used in this investigation. Plateau voltage curves were run on the detector in conjunction with varying gain settings in order to determine the peak operating voltage of the de- tector. The pulse height selector settings were set exactly at the base of the MoKa energy peak. Table 2 Machine Parameters SLIT SIZES Front Entrance 100 p Rear Exit 200 p x 1 cm SPECIMEN TO DETECTOR DISTANCE 215 mm OPERATING CONDITIONS 45 KV 20 ma o RADIATION MoKa X = .71 A DETECTOR VOLTAGE 1100 V (Bicron* scintillation) PULSE HEIGHT SELECTOR Lower Limit 2.0 volts Upper Limit 4.0 volts STEP SIZE 100p *Bicron Corporation, Newbury, Ohio All samples were step scanned using an Ortec* Model 6713 axis controller and Ortec counting equipment. Scans began from a height of approximately 300 p above the cen- troid of the primary beam to about 10,000 p above the main beam. Collection time was 1,000 seconds or at least 12,000 counts at each of the 100 p steps. All scans were run with the Kratky block, specimen holder and tank evac- uated in order to reduce air scatter. WOrtec, Inc., Oak Ridge, Tennessee 62 A sample of Lockheed pyrolitic graphite was used both as a test of proper camera alignment and as a scat- tering standard to correct for variations in primary beam intensity between runs. Figure 8ashows the scattering curve of the graphite, while 8b shows the Porod region, where a slope of -3 indicates proper camera alignment. All intensity corrections were made with the graphite standard. Main beam intensity variations were corrected for by ratioing the intensity of the graphite scattered at a fixed position before each run to that of the graphite scatterer measured before the pure aluminum run. Varia- tions in sample absorption were corrected for by measuring the intensity of the graphite standard with the pure alu- minum in the absorption position and ratioing it to that of the graphite standard with the aluminum alloy sample in the absorption position. The total corrected intensity for a sample is given by: istd Apure 1 = abs corr meas Isam sam g abs where: Icorr is the corrected intensity Imeas is the measured intensity I s is the graphite intensity measured prior to g each run a) rI co 0 C) P4 4-i 0 a) a) 0 a) C-3 4-3 W co a0 rZ4 60 I CD co I v rn I / 3 Q i I 0 / (- i 0 C) SI , / I o 0 2 o C a LLa N LN ~5/'3 AIS 5rN 61 GRAPHITE /00 O\ -\ 13 90- 80- hF 70 - Z 60- .- 50 - 40- 30 40 50 60 70 80 90 6 (RFADIAN xl -2 Figure 8b. Porod Region of Lockheed Pyrolitic Graphite Scatterer std Istd is the graphite intensity measured prior to g the pure aluminum run Apure is the absorption factor of the pure alu- abs minum Asam is the absorption factor of the sample abs The sample scattering curves were corrected for back- ground by subtracting the intensity scattered from a 99.999% pure aluminum sample with the proper correction factors described above applied. All data was plotted and smoothed by hand before evaluation by a computer pro- gram written by Gould and Kirkli,6 and modified by Healey and Hill64 for this work. This program corrects for pri- mary beam intensity variation due to sample thickness and tube fluctuations automatically, and calculates the Guinier Radius, Porod Radius, integrated intensity and the particle size distribution, based upon the method of Harkness, Gould and Hren. 50 Resistometry Measurements Resistometry measurements were made using a Leeds and Northrup* Kelvin bridge-type potentiometer, standardized 65 against a .001 ohm standard resistor. Measurements were made at liquid nitrogen temperature. It was not possible *Leeds and Northrup Company, Philadelphia, Pennsylvania to measure true resistivity (p) of the sample as the exact length between the voltage leads was not known. Instead, the resistance, or actually the voltage between the two welded leads with an exact current of 1 amp, was measured. R-R This was plotted as R where Ro was the resistance o measured immediately after quenching. Tensile Testing Tensile tests were performed on the specimens cut to ASTM Specification A370 (substandard size) on an Instron* testing matching. All tests were run at a strain rate of .02 inches per minute. *Instron Corporation, Canton, Massachusetts CHAPTER IV EXPERIMENTAL RESULTS All data presented here has been obtained from com- mercial 7075 alloy, whose chemical composition is pre- sented elsewhere. A series of isothermal aging treat- ments, with and without deformation, were performed on the samples. The term preaging, as used here, will include any thermal treatment given to a series of samples prior to the final thermal aging treatment. Table 3 presents all treatments performed on the various series and their designations. Resistivity The resistivity-aging time curves for the wire speci- mens are given in Figure 9. This resistivity data is sum- marized in Figure 10 as an Arrhenius plot.33 The physi- cal significance of this plot lies in the slope of the line having the units of energy, and corresponding to Em, the activation energy of motion. 0 0 0 C 0 0 0 0 0 0 0 JI 0 Si 09 0,( 0 0 0p IC~ / I. i I i I I l I i. 'c aCC d 000 o 0 6!--U b 0 0 0 0 0 0- S I* I* Ic 0 0u0 .o *e e I') (/ -! -- 100,000 - 90,000- 80,000 - 70,000- 60,000 - 50,000 - S40,000- 30,000- S20,000 - 10,000 9,000 8,000 7,000 6,000 5,000 4,000 - 3,000 - 2,000)- I I 2.60 2.55 2.40 100o/T Figure 10. Arrhenius Plot of Resistivity Maxima 2.75 2.70 2.65 2.50 2.45 , WI .V 1 I rnni Table 3 Summary of Aging Series Preage Treatment Aging NONE NONE NONE NONE NONE 1 hour 125C NONE NONE 100 hours 750C 25 hours 75C 100 hours 95C 25 hours 950C 5 hours 125C Temperature 1350C 1350C 1350C 125C 1250C 1250C 750C 75C 1250C 1250C 1550C 1550C Small Angle X-Ray Scattering Preliminary Studies of Initial Aging A similar study to that of resistivity was performed using small angle scattering. A single sample was iso- thermally aged at three temperatures (1350C, 1450C and 1650C) with homogenization at 4600C for two hours between temperatures, and the scattered intensity was measured at Series A B D E F F1H G H J K L M Defm 0 5% 15% 0 15% 15% 0 15% 0 0 0 0 0.011 radians as a function of aging time. Figure 11 shows the scattering curves obtained, while Figure 12 is an Arrhenius plot of the log of the time to reach a scat- tering intensity of 1.5 counts per second versus 1000/Ta, where Ta is the isothermal aging temperature.40 A preliminary study of the effects of deformation on the initial aging rate was performed in a similar manner. The scattering intensity at 0.011 radians is plotted ver- sus aging time for the A, B and D series, see Figure 13. Change in Zone Size Parameter Entire scattering curves are necessary to determine the two size parameters; the Guinier Radius, RG; and, the Porod Radius, R Figures 14 and 15 show two such scat- tering curves, one representative of a direct-aged sample (Figure 14) and the other a sample that had been preaged before the final aging treatment (Figure 15). Corrections for sample variation were performed according to the pro- cedure presented in Chapter III. The results of the evol- ution of the Guinier Radius are given in Figures 16 through 21. Evaluation of the Guinier Radius in the unaged or very short direct aged samples was difficult due to the low scattering intensity. For the low temperature aging series (G and H), this problem was encountered throughout the aging sequence as the growth was extremely slow. 69 \~g Or o0 r--I r--4 * (Q r LU C)2 \ x (0 ,Q -o D .CL ~- > m *\3 a 0 \ 0 r-4 II) o m :l I \IO I" \ \ <. \ v .50 2.40 2.30 2.20 100/A Figure 12. Arrhenius Plot of Time to Reach Scattered Intensity of 1.5 Counts Per Second 20,000- /0,000 9,000 8,000 7,000 6,COO 5,000 4,000 3,000 2,000 1,0001 2 *k 4-i CH 41 1-4 41i 4i Ci ca Q) W C) P4 0 0 C) 0 41i 0 -e 4-i 44 *r-4 r4i C) Cf) bl P4 -I @4K Nl S 4. Cl) :1 I. (LD 0- SNVI/ V/dI I/10 v C(Sdc) AIISN31NI ro rr> c) rE 101 4i (0 bc u 4,-i U) cri bo 04 Fh 74 .0 I(C) Fc)c rnr ~t I LIt- // .7 UI I (SdD) AIISN31NI a) a) a) cli C- 0 a) U cli r-4 a) r4J C-q cli cl .r4 a) S a) r-4 Lr)L 76 Lii O I * S-I -- Q ! So I r) -- .---7 SI I I I (SdJ) AIISN31NI Ii SII SII H1 I I * U S QQQ 00 o S. - snRaVd 31N/RND o0 ro 0 ) ^ 0 I- _ IY~-LU-_ ~~ I I0 cn 0) r4 1144 14 0 0) r4 0) 4- 0 .r-4 0) r4 *1- 0 L4a 0, 0 *1- '-4 0 0) I r \ 4\ \ I II I I Q CO co I1 Ic LUKK^ (1) Q) KI) i Cr 0 V &FIlG' I 31iN1,A9 0 1 U0 ~44 ~ 44 Z - 02 w w (3) 0 co p 0 44 rz a) 0 CO j 10 a) Cd 44; 0 r. :j -r4 rl4 Lo 2t: 4 u U U SS I'C v s'f/CVU' UJiV/fl c$ U) .r4 0 LI:4 $4 a) -,4 4-4 0 r- LIe4 0 0 41 0 cii ON rL, *1 i 83 o 0 0 0 o N 00 > o L C L N C- 0 Cf) \ f)3 UNS d'A \. I2S n \, ^^^ \ * -^ \- x s \ \ | i 84 o So 0 cn X Ci Q IC) - ;J O0 O O " O o O .-1 Q (I) ar ,. o- \) t .- I * \ 0o \ 4-'0 \0L \ ^ '5 \ -^R O ; <-u \ ^ *ri V) a) 0 a) rn O c) 0 0 4.r 41 0 a) C4 :4 rl 86 0 0 0 oco 00 0 L.) crc QQ 0 o o ias -a , co 0 P^ ) co to v o <0 < kj
50 The results for the Porod radius are presented in Figures 22 through 26. It was not possible to obtain a meaningful Porod radius for the very short aging time samples as the Porod region was either at too high an angle, or buried within the background. Also, it was not possible to obtain a Porod radius with physical signifi- cance for the G and H series as the Porod region was be- yond the workable angular range of the Kratky camera. The radius of maximum abundance values are plotted in Figures 27 through 32. These determinations assume a log normal distribution. It was not possible to obtain a radius of maximum abundance for samples for which no Porod radius could be calculated. The Porod radius is necessary for this deter- mination as described in Chapter II. Frequency distribution plots of the particle radii as a function of aging time are presented in Figures 33 through 41. These distributions were generated assuming a log normal distribution, and spherical particles, and again were generated only for samples for which it was possible to obtain both a Guinier and Porod radius. All of the above parameters were calculated using the modified computer program of Gould and Kirklin63-64 which is described in the Appendix. |