Interaction of dislocations and structures in cyclically strained aluminum alloys

INTERACTION OF DISLOCATIONS AND STRUCTURES IN CYCLICALLY STRAINED ALUMINUM ALLOYS*

J. B. CLARK? and A. J. McEVILYt

The effect of cyclic straining on the microstructure of Al-4 wt. % Cu and Al-15 wt. % Ag alloys in the solution treated and aged conditions has been studied by electron transmission microscopy. During the cyclic straining of these alloys, the copper and silver solute atoms limit the substructure formation and alter the nature of the loops generated in contrast to the complex dislocation tangles and well-developed cell structure produced in pure aluminum. The interaction of the dislocations produced by cyclic straining with the metestable hardening structures in these age hardening alloys can generate soft regions in the microstructure where the probability of crack nucleation will be increased. In the Al-4 wt. x Cu, the G.P. II zones undergo reversion whereas in Al-15 wt. % Ag, the nucleation and growth of transition structures with concomitant dissolution of G.P. zones is promoted by the cyclic straining. It is concluded that the disproportionately low ratio of endurance limit to tensile strength of the high strength aluminum alloys probably is due to the instability of the strengthening structures of the alloys during cyclic straining.

INTERACTION DES DISLOCATIONS ET STRUCTURES DANS DES ALLIAGES D’ALUMINIUM SOUMIS A UNE DEFORMATION CYCLIQUE

L’effet d’une deformation oyclique sur la microstructure des alliages Al-4 % Cu et Al-15 % Ag, mis en solution et vieillis, a fait I’objet d’une etude en microscopic Blectronique par transmission. Pendant la deformation cyclique de ces alliages, les atomes de cuivre et d’argent en solution, limitent la formation des sous-structures et alterant la nature des boucles de dislocations qui apparaissent, en comparaison aux Bcheveaux complexes de dislocations et 8. la structure cellullaire bien developpee que l’on observe dans l’aluminium pur. L’interaction des dislocations produite par la deformation cyclique, avec les structures m&astables responsables du durcissement dans ces alliages vieillis, peut provoquer l’apparition de regions non-durcies dans la microstructure aux endroits oh la probabilite de formation de fissures sera augment&e.

Dans l’alliage Al-4% Cu, les zones de Guinier-Preston II, subissent une reversion tandis que dens l’alliage Al-15% Ag, la nucleation et la croissance de structures de transition avec la dissolution simultanee de zones de Guinier-Preston est favorisee par la deformation cyclique. En conclusion, la faible valeur du rapport entre la charge de rupture des alliages d’aluminium L haute resistance est probablement due a I’instabilite des structures responsables du durcissement des alliages au tours de la deformation cyclique.

WECHSELWIRKUNG VON VERSETZUNGEN UND AUSHARTUNGS-STRUKTUREN IN ZYKLISCH VERFORMTEN ALUMINIUM-LEGIERUNGEN

Mit Hilfe von Elektronen-Durchstrahlungs-Mikroskopie wurde der EinfluLi zyklischer Verformung auf die Mikrostruktur der Legierungen Al-4Gew. %Cu und Al-15Gew. %Ag bestimmt, und zwar im Zustand nach dem LBsungsghihen und nach Auslagerung. WSlhrend der zyklischen Verformung dieser Legierungen schranken die gel&ten Kupfer- und Silberatome die Bildung von Substrukturen sin, so da0 die. Art der erzeugten Versetzungsschleifen sich von den komplizierten Versetzungsknaueln und der wohlentwickelten Zellstruktur des reinen Aluminiums unterscheidet. Die Wechselwirkung der durch zyklische Verformung hervorgerufenen Versetzungen mit den metastabilen Aushartungs-Strukturen dieser aushiirtbaren Legierungen kann zur Entstehung weicher Bereiche in der Mikrostruktur ftihren, in denen die Wahrscheinlichkeit zur Bildung von RiBkeimen erhijht ist. In der Legierung Al-4Gew. %cU tritt Riickbildung der G.P.II-Zonen ein, w&rend die zyklisohe Verformung bei Al-16Gew. %Ag dazu ftihrt, daD Keime von Vbergangsstrukturen gebildet werden und unter gleichzeitiger Aufliisung von G.P.-Zonen wachsen. Es wird geschlossen, dell das ungewijhnlich niedrige Verhiiltnis von Dauerfestigkeit zur Zugfestigkeit hochfester Aluminium-Legierungen mit der Instabilitat der Hartungsstrukturen dieser Legierungen bei zyklischer Verformung zusammenhangt.

INTRODUCTION

The nature of dislocation distributions in cyclically loaded pure metals has recently been studied by the use of electron transmission thin film microscopy. For example, a substructure bounded by a three- dimensional network of tangled dislocations has been observed in aluminum when strained above a certain strain amplitude. (1) Other work has shown that complex dislocation tangles as well as dislocation dipoles are formed in topper,(2) aluminum,(1~3*4) mag-

* Received April 22, 1964. t Scientific Laboratory, Ford Motor Company, Dearborn,

Michigan.

nesium and niobium.(s) and a-brass.@) Such direct observations contribute significantly to our knowledge of the processes taking place in cyclic deformation. However there has been little information available on the effect of cyclic loading on dislocation structure in multi-phase systems, in which interactions with dispersed second phases are involved.

Age hardening aluminum alloys are particularly interesting because evidence in the literature indicates that the metastable structures responsible for the strength of these alloys undergo a transformation that leads to localized softening during cyclic loading.(73) Such behavior has been attributed to either overaging

ACTA METALLURGICA, VOL. 12, DECEMBER 1964 1369

4

1360 ACTA METALLURGICA, VOL. 12, 1964

TABLE 1

Alloy Heat treatment structur0

Al-4 wt. % cu

ill-l.5 wt. O/, Ae

Solution treatment 16 hr at 540°C + plate quench ST + 16 hr at 130°C ST + 5 hr at 160°C ST + 2-l/2 hrs at 250°C Solution treatment

OL solid solution tc + G.P. I zones CL + G.P. II zones CL + 0’ transition lrtttice

16 hr at 540°C + plate quench ST + 120 hr at 160°C

CL solid solution a + G.P. zones + y’ transition lattice

or reversion of the mete&able structure, but direct evidence for either process has been lacking.

Therefore, this electron microscopy study was undertaken in order to investigate more fully the nature of dislocation interactions with typical structures under cyclic loading and to obtain direct evidence of the cause and nature of the instability process in age hardened aluminum alloys.

MATERIALS AND TESTS

Two alloys (Al4 wt.% Cu and Al-15 wt.% Ag) were selected for investigation because of the detailed knowledge of the precipitation processes in these alloys and because a variety of microstructures can be readily obtained for the study of ~slo~tion- precipitate interactions during cyclic loading. The super-saturated solid solution of the Al4 wt.% Cu alloy rejects excess copper in more than one form according to the time and temperature of aging, namely: G.P. I zones, G.P. II zones (or 8”),* a transition lattice 8’ and the equilibrium 8 (C&&l,) structure. In the Al-15 wt.% Ag alloy the forms sre G.P. zones, IG transition lattice y’, and the equilibrium y fAgsAl) structure. In the Al-Ag system the G.P. zones are spherical, in contrast to the disc shaped G.P. zones of the Al-Cu alloy. The microstructures selected for study and the concomitant heat treatments are listed in Table 1.

Most of the thin film specimens utilized in this investigation were obtained from 0.010 in. sheet. After heat treatment, specimens from this thin sheet were glued to an &luminum alloy subst~~ and cyclically strained in fully reversed torsion for 105 cycles at a shear strain amplitude of 0.002 at a rate of 1800 c/mm. After cyclic loading, the specimens were removed from the substrate and thinned to 1000-2500 A by the window technique using Lenoir’s elect~polishing sol~~tion.~ll) The thinned foil was

* This structure has been named G.P. II by Silcock et aZ.‘g’ More recently, Nicholson et ol.(r”i have suggested the symbol 8” since they feel that this phase has a. de&&e crystal stratus and is not a simple perturbation of the matrix structure. However in the opinion of the authors, the evidence is still insufficient to definitely classify this structure as a transition lattice which the symbol 8” implies; therefore the older symbol G.P. II is used in this p&per.

cleaned in a special stripping solution(12) to remove electropolishing films and examined in an Hitachi HU- 11 microscope operated at 106 kV.

Fatigue bars, 0.06 in. thick, of the alloys in the various heat treated conditions were subjected to cyclic loading to the same strain amplitude es the 0.010 in. sheet samples to ascertain whether the observed dislocation arrays depended upon specimen thickness. The fatigued sress of the bars were reduced from one side to a thickness of 0.025 in. with a Servo Met Spark Cutter. They were then eIectro- polished on one side and thinned toward the outer surface until reduced to 0.010 in. thick followed by electropolishing on both sides to final foil thickness. For a given thermo-mechanical history, similar structures were observed in both types of fatigue specimens.

Since the reduction of a specimen to a thin foil for examination by electron transmission prevents examination of the Same region of a specimen before and after fatigue, representative sampling of the fatigue- induced structures is necessary. However, a sampling problem arises in the application of the electron transmission technique to the study of fatigue deformation because the region examined in a given specimen by this method is necessarily smrtll and fatigue deformation is a highly localized process. For this reason, for a given alloy microst~ctu~ and degree of cyclic strain, many foils were prepared and examined. Only consistent observations based on a number of specimens were considered to be truly representative of the dislocation-precipitation interactions for the given conditions.

Al-4 WT.% Cu ALLOY

Solution treated structure. Examination of thin films of the solution treated and quenched materid which had not been cyclically loaded revealed the presence of a few dislocation loops formed as a result of vacancy condensation. In addition, long helical dislocations were observed. These helical dislocations form during quenching of the alloy as a result of climb due to the interaction of vacancies with dislocations of strong

CLARK AND McEVILY: CYCLICALLY STRAINED ALUMISUM 1361

FIG. 1. Al-4 wt. % Cu. Solution treated + 100 k cycles at *0.0017 shear strain. Loops and helical dislocations produced by cyclic straining. x 40,000

screw character.cf3) After cyclic loading, the helical dislocations were distorted and the number of loops was increased as shown in Figs. l-3. Similar effects have been observed by Swanno4) for the case of static unidirectional loading of an Al4 wt.% Cu alloy.

The character of the loops produced in the solution treated alloy contrasts greatly with that seen after cyclic straining of pure aluminum in which a definite cell structure and a large number of dipole loops elongated in the (112) directions are generated.(l) Also no cell structure of dislocation tangles was observed in the solution treated alloys.

These differences between the b-4 wt. % Cu alloy and pure aluminum appear to be due to a lower stacking fault energy of the alloy with respect to aluminum. Such a reduction would inhibit the cross- slip required for dipole formation and for the formation of a cell structure. It is noted that the addition of magnesium to aluminum also tends to inhibit the formation of a cell structure.(15)

G.P. I structul-e. If the Al4 wt. % Cu alloy is aged for 16 hr at 13O”C, small G.P. I zones of copper with disc shape (~3-6 A thick and 80 A in die.) are formed with an average separation of 100 A.(la) These G.P. I zones are weak and as reported by Kelly and Fine(“) such zones are easily cut by dislocations.

Cyclic loading of alloys containing G.P. I produces a large number of uniformly distributed dislocations

(Fig. 4). Because the zones are very small, and show only as a mottling of the background in Fig. 4, it is difficult to identify specifically the dislocation-G.P. I zone interactions. The dislocations are slightly bowed suggesting that the dislocations are impeded but not arrested by these weak zones. In a few foils, faint structures in the background suggested the formation of zone-free regions induced by the cyclic loading but the smallness of the G.P. I zones pre- cluded definitive evidence. However since the G.P. I zones are weak and easily cut, interaction of the dislocations with the zones to form preferentially weak regions in the structure is not as important as in the case of stronger G.P. II zones discussed below. No evidence of cell formation, or precipitation of G.P. II zones or of the 8’ transition lattice, indicative of overaging, was observed.

G.P. II structure. If the A1-4 wt.% Cu alloy is aged 5 hr at 160°C large disc shaped G.P. II zenes (~20-30 A thick and 400 A in dia.)(i@ form and maximum hardness is attained. During deformation, dislocations cut these G.P. II zones because the average separation of the zones is only 300 A(ls) and it is energetically easier to cut the zones than to bend the dislocations between the pinning poinWa) or climb around them. A characteristic feature of cyclic loading of this microstructure was the localization of dislocation motion to a greater degree than in any of the other microstructures studied. Deformation


FIQ . 2. Al- -4 wt. % Cu. Solution treated + 100 k cycles at +0.0017 shear strain. Helical dislocations loops in [loll directions. Note the type of contrast around the loops. x 40,000

an d rows of

FIG. 3. Al-4 wt. % Cu. Solution treated + 100 k cycles et I_tO.O017 shear strain. Helical dislocations produced by cyclic straining. Note the interactions of the dislocations and the irregularity of the helices. x 40,000

CLARK AND McEVILY: CYCLICALLY STRAINED ALUMINUM

FIG. 4. 41-4 wt.% Cu. Solution treated + 16 hr at 130°C + 100 k cycles et 10.0017 shear strain. Uniformly distributed dislocations cutting G.P. I zones. x 30,000

appears to be confined in regions where localization of strain end dislocation sources promote shearing of the G.P. II zones. After cyclic loading of G.P. II microstructure, precipitate-free bends, about 1.5 p wide, filled with dislocation arrays were observed, as shown in Fig. 5. In these bands, the hardening precipitate had disappeared. Apparently with concentration of the strain, the passage of a large number of dislocations through a limited number of G.P. II zones fragments the zones to a size that offers little or no resistance to further dislocation motion. The zones are so broken up that effective reversion of the zone has occurred. Recently, Forsyth(lg) has reported the observation of similar precipitate-free bands in an Al-Mg-Zn alloy.

Efforts to detect the presence of extremely small G.P. II zones within these soft regions by direct observation in the electron microscope were un- successful, and it is therefore uncertain whether the G.P. II zones in these precipitate-free regions have been fragmented to a size below the resolution of the microscope or have actually been reduced to a critical size and reverted to form a region of supersaturated solid solution. However, the large size of the G.P. zones bordering the precipitate-free regions suggests that these zones had grown into a supersaturated solid solution and that reversion of G.P. II zones had in fact occurred in this alloy. These results support

the prior evidence(zo) for reversion during fatigue of 2024-T4, an Al-Cu base alloy, which also hardens by precipitation of G.P. zones.

Experiments were made in which bulk specimens cycled to develop these reverted regions were re- annealed at 160°C to see if the presence of the large number of dislocations which might act as nucleating sites would influence the reprecipitation process in these zones. The complete absence of reverted zones in these specimens indicated that the heat treatment eliminated the soft regions by reprecipitation of G.P. II zones in normal size and dispersion. Precipitation of 0’ by a reanneal at 160°C was not facilitated by the presence of the dislocations induced by cyclic straining. Nucleation of 8’ on these dislocations is a difficult process. The temperature of the reanneal had to be raised to at least 200°C before 0’ nucleated on the dislocations in the zone-free regions. The pertinence of this difficulty of nucleation of 8 with respect to the formation of these soft regions free of G.P. II zones will be discussed below.

8’ structure. If the Al4 wt. % Cu alloy is aged for 2.5 hr at 250°C to produce platelets of the semi- coherent transition lattice, W, in a matrix devoid of G.P. zones, then, during cyclic loading, extensive dislocation motion occurs in the depleted aluminum matrix. Evidence for this is given by the large number of tangled dislocations illustrated in Fig. 6. The


FIQ. 5. Al-4 wt. ok Cu. Solution treated + 5 hr at 16O’C + 100 kcyoles at rtO.0017 shear strain. Precipitate- free region produced by cyclic loading. Note the arrays of dislocations in the precipitate-free region 8nd also that

the G.P. II zones are larger along the edge of the region of apparent solid solution. x 20,000

FIG. 6. Al-4 wt. % Cu. Solution treated + 2.5 hr at 250°C + 100 k cycles. Dislocation loops at 8’ platelets and complex dislocation tangles in dilute alloy matrix. x 15,000

CLARK AND McEVILY: CYCLICALLY STRAINED ALUMINUM 1365

FIG. 7. Al-4 wt.‘% Cu. Solution treated + 2.5 hr at 250°C + 100 k cycles at &to.0014 shear strain. Complex interactions between dislocations and 8 platelets: et A, a dipole is forming; at B, interaction at the W/matrix

interface; at C, a loop around a 8’ platelet is evident. x 30,000

fuzziness around the 0’ platelets is due to dislocation loops being truncated during thinning of the foil. Superficially, the dislocation interactions are similar to those reported by Koda et uZ.(~~) and Bonar and Kelly.(22) However, as shown in Figs. 6-9, closer examination indicates that the reversed straining produced more complex tangles. In passing it might be of interest to note in Fig. 7 that dipole formation is shown at A and loop formation is seen at C. With a slight tilt of this field, slip traces are in contrast and show the presence of cross clip around the 8’ platelet. At B, the change in the fringe pattern of the W/matrix interface denotes the interaction of the dislocations with this interface and possibly shearing of the 13’ platelets. However, such interactions are rare and the general fine striations on the platelet interface, noted by Koda et uZ.(~~) for unidirectional deformation, which they interpreted as fine slip in the 6’ platelets, was not seen after cyclic loading of the same structure in the present study. As shown in Fig. 8, the cyclic loading produces complex tangles rather than the straight uniform arrays seen after unidirectional deformation.

Gross shearing and bending of the 8’ platelets was never observed after cyclic straining in agreement

with results of Bonar and Kelly(22) for unidirectional deformation of this structure at small or moderate strains. Even in the vicinity of a crack when complex tangles and dipole loops form, as shown in Fig. 9, deformation of the semi-coherent 8’ precipitate does not occur during cyclic loading at these shear strain (fO.002) levels. The interface patterns seen in Fig. 9 are Moire patterns and not fine slip striations as proposed by Koda.t21) Thus, in this microstructure the interaction of the 8’ platelets with the dislocations acts to promote dislocation tangles in the matrix, rather than to create weak regions by altering the 8’ structure.

Price and Kelly (23) have shown that in an Al-Cu alloy with the 8’ microstructure, the yield strength is relatively low and the work hardening rate is very high due to the high density of dislocation loops at the 8’ particles. The dislocation arrays in the matrix, shown in Figs. 6 and 9, resemble more closely that observed in pure aluminum rather than that observed in the solution treated alloy. The fact that dipole loops and cell structure of dislocation tangles are seen in the dilute matrix after cyclic loading indicates that the matrix solute is appreciably reduced. With the presence of these soft regions in the microstructure


FIG. int

8. erf

Al-4 wt. % Cu. Solution treated + 2.5 hr at 250°C + 100 k cycles at &to.0014 shear strain. ice tangles and the bending of the dislocations around 0’ platelets. Plane of foil (100).

N .ote the X1 10,000

FICA 9, *a

Al-4 wt. % Cu. Solution treated + 2.5 hr at 250°C + 100 k cycles at &IO.0014 shear strain. fracture. Note the complex tangles and dipole loops in the matrix without gross deformat

platelets. Moire fringes are seen in sorne of the 13’ platelets. x 30,000

. R ,ion

legion of e

CLARK AND McEVILY: CYCLICALLY STRAINED ALUMIKUM 1367

FIG. 10. Al-15 wt. ‘A Ag. Solution treated + 100 kc/s at -&O.OOll shear stain. Loops and helical dislocations produced during cyclic loading. x 32,000

no breakdown of the 8’ platelets is required to initiate fatigue failure.

Al-15 WT.% Ag ALLOY

~oZz&o~ treated structure. As shown in Fig. 10, cyclic loading of the Al-15 wt. % Ag alloy in the solution treated condition produces a large number of loops essentially similar to those observed in solution treated Al4 wt. % Cu alloy. As in the _A-Cu alloy, no cell structure of dislocation tangles was seen in this alloy.

Aged ~t~~~~ (G.P. Zones + y’). For the Al-15 wt. *A Ag, the aging treatment used by Nicholson and Nutting’24) i.e. 120 hr at 16O*C, produces the microstructure shown in Fig. 11, which consists of a few pl&elets of the transition structure y’ and large (90 A) spherical G.P. zones. Kate the zone-free regions adjacent to the y’ platelets. Some discussion on the formation of this microstructure may be appropriate in order to comprehend the effects of cyclic loading.

The G.P. zones are known to nucleate homo- geneously,(25?26) possibly by a spinodal reaction, within a few minutes at 160°C.(e7) The y’ phase however requires heterogeneous nucleation. After a standard solution treatment and water quench approximately 20 hr of aging at 160” are required before y’ can be detected by X-rays.(27) The y’

platelets nucleate heterogeneously at isolated stacking faults associated with sessile dislocations formed during quenching, at helical dislocations@4) and at ordinary dislocations. t2*) Silver collecting at these sites lowers the stacking fault energy. Growth of the y’ parallel to its basal plane occurs by dissociation of the dislocation and propagation of the interfacial partial dislocations at a rate dependent on the rate of diffusion of silver atoms.@*) The fact that dislocations facilitate the heterogeneous precipitation of y’ by seting as nucleation sites is shown in Fig. 12. In this case the alloy was cold rolled 50% after solution treatment and aged only 24 hr at 160°C. The much larger proportion of y’ is evident. Thus the paucity of y’ platelets formed after 120 hr at 160°, shown in Fig. 11, results from difficulty of nucleation rather than slow growth.

During aging the y’ platelets, having a lower energy than the G.P zones, will continue to grow and to deplete the surrounding matrix. When the solute concentration of the matrix, near the y’ platelets, is reduced below that of the immiscibility gap for the G.P. zones,(25f the G.P. zones will dissolve, leaving a zone-free region adjacent to the y’ platelets 8s shown in Fig. 11. Hren and Thomas(2*) have shown that during continual aging these zone-free regions continue to widen as the y’ platelets grow until eventually all the G.P. zones disappear, At 160°C and below, the


FIG. 11. Al-15 wt. % Ag. Solution treated + 120 hr at 160°C. Platelets of the transition lattice y’ and large spherical silver rich G.P. zones. Note the zone-free regions adjacent to the y’ platelets. x 30,000

FIG. 12. Al-15 wt. % Ag. Solution treated + cold rolled 50 ‘A + 24 hr at 160°C. Note the much larger proportion of y’ due to the increased nucleation from cold rolling prior to aging. x 20,000


FIG. 13. Al-15 wt. ‘4 Ag. Solution treated + 120 hr at 160°C + 400 kcycles at *0.0012 shear strain. X’ote the increased amount of y or y’ in [loll directions after cyclic straining as compared with the initial amount of y’

shown in Fig. 11 prior to straining.

supersaturation of these zone-free regions is very small because at temperatures below 200°C the lower limit of the immiscibility gap for the zones lies very close (<0.25 wt. % Ag) to the solvi of the equilibrium phases.(25) Since the solvus of transition lattice would lie even closer to the immiscibility gap than the equilibrium solvus, the supersaturation of these zone-free regions must be essentially zero. Obviously these depleted regions will be softer than the surrounding zone-filled matrix.

It should be expected then in considering the effect of cyclic loading that the main resistance to dislocation motion in the structure shown in Fig. 11 will be the large (90 A) closely spaced spherical G.P. zones. The relatively small amount of y’ present after 120 hr at 160°C will play a small role in the hardening of the alloy.

The effect of cyclic loading on the Al-15 wt. % Ag with the initial microstructure shown in Fig. 11 is given in Fig. 13. The amount of y or y’ has increased markedly and is preferentially oriented in the [loll directions. Therefore cyclic loading tends to cause overaging in an A.-15 wt. ‘A Ag alloy containing the initial microstructure as shown in Fig. 11. The term “over-aging” is used here to denote the formation of the transition lattice y’ and/or its transformation to the equilibrium precipitate y. The stimulation of these overaging reactions by cyclic loading was con-

firmed by X-ray diffractometer studies of cyclically strained and unstrained test bars of the alloy using a General Electric diffractometer with cobalt radiation and curved lithium fluoride monochromator. Despite the preferred orientation of the bars and the small amounts of y’ or y, the (lOi1) peaks of these phases could be measured when either individually was present. Unfortunately the peaks overlap, so that when both are present, it is not possible to determine explicitly the ratios of y’/y as a function of cyclic loading. However as cyclic loading progressed, the size of (IOil) peak of the y’ increased and moved toward that of equilibrium structure y, indicating the increased formation of y’ and its transformation to the equilibrium y phase. It is concluded that overaging rather than reversion occurs in the Al-15 wt. ‘A Ag alloy during cyclic loading.

Cyclic loading induces overaging in this alloy by providing dislocations required for the nucleation of the y’.t2s) Oscillating dislocations with jogs, moving non-conservatively, will increase the number of vacancies and thereby locally accelerate the growth of the y’ platelets. Adair et a1.(2g) have shown that the creation of vacancies by radiation can produce rapid overaging of aluminum age hardening alloys at low temperatures. Careful examination after cyclic loading shows a build-up of loops around the y’ platelets. At a critical time in the growth of the y’ platelets, it


is possible that these dislocations could initiate the transformation of y’ to y by stimulating the breakdown of the coherent ~‘~rnat~x interface and the production of the dislocation boundary between the equilibrium phase y and the matrix. This instability of the y’ with respect to deformation has been confirmed recently by Laird@O) who found that the filing of an Al-15 wt. % Ag alloy consisting initially of only G.P. zones +y’, effectively converts all the y’ to y. Thus the conversion of y’ to y during cyclic straining may be expected.

As shown in Fig. 11, the growth of the y’ platelets depletes the matrix locally and causes the G.P zones to dissolve leaving soft precipitate free regions in the matrix. Conceivably, in regions of high deformation, the fatigue-induced growth of y’ could continue with concomitant dissolution of G.P. zones, thus producing in local regions, miorostructures and dislocation interactions similar to those reported for the 8’ platelets for the Al-Cu alloy shown in Figs. 6-9. However in this study such regions in which fatigue induced overaging had progressed to such an ad- vanced stage were not actually observed, though the evidence that the Al-15 wt. % Ag alloy with the initial microstructure shown in Fig. 11 does overage during cyclic loading is conclusive.

DISCUSSION

It is seen that in certain microstructures in these age hardening alloys, cyclic straining produces soft regions in which further deformation is localized and the probability of crack nucleation significantly increased. The weak regions in the Al4 wt.% Cu alloy are produced by cyclic strain-induced reversion of the G.P. II zones, and in the A-1-15 wt. % Ag alloy by dissolution of the G.P. zones as a result of cyclic strain-educed overaging reaction (formation of y’ and y). Thus the hypotheses of Broom et aZ.(7>8) that the metastable hardening precipitates in aluminum age hardening alloys are unstable during cyclic loading are confirmed. This basic response of these alloys to cyclic straining is responsible for the low fatigue properties exhibited by the~ghstrength agehardening alloys. In the AI-Cu alloys, the very structure that produces the high strength, the G.P. II zones, undergoes reversion during cyclic straining. Probably some microstructures of the commercial aluminum alloys will react to cyclic Ioading, in the manner of the Al-Ag alloy and overage in local regions. There- fore in designing aluminum alloys, for cyclic loading, the stability of the hardening structure must be considered because, as in creep, thermo-kinetic processes can and do occur during cyclic loading

despite the lower temperatures at which cyclic loading is normally imposed.

With certain aged structures in the Al-Cu alloys, cyclic straining does not need to alter the microstructures for failure to occur, since at least, part of the microstructure is weak and thus deformation with subsequent crack nucleation occurs readily. For example, the G.P. I zones are weak and easily cut. Crack nucleation will occur by normal mechanisms and any alteration of the G.P. I zones by cyclic deformation probably has little influence on the failure process. Similarly, in the case of the 8’ platelets in a matrix devoid of G.P. zones, deformation occurs readily in the weak matrix. Alteration of 8’ was not observed and is not a requisite because the matrix is intrinsically weak as a result of solute depletion and thus cracks can readily nucleate.

The question arises as to why cyclic straining produces reversion of the G.P. II structure of the Al4 wt. % Cu alloy and overaging reactions in the G.P. zone + y’ structure of the Al-15 wt. % Ag alloy. In both cases, dislocations cut the G.P. zones. The cutting of the solute-solute bonds is, in effect, the first stage of reversion process. Why then the difference in the behavior of the two alloys1

The answer lies in the relative rates of nucleation of the transition lattices, 0’ and y’, of the two alloy systems. When the zones are cut by dislocations, the energy level of the zones is raised due to the destruction of bonds and the increase in surface area. The zone may ~ont~ue to be cut to finer and finer size due to the large number of cutting dislocations coursing back and forth during cyclic straining leading eventually to reversion of the zones and resulting in localized softening. The relatively high strength of the adjacent precipitation hardened regions will tend to confine the dislocation motion to the narrow reverted band. This would be con-

sistent with the observations that fatigue damage is difficult to detect in age hardened aluminum alloys prior to crack initiation. A fatigue crack at the surface could be relatively easily initiated in the softened region. On the other hand, if the cyclic loading induces competitive overaging reactions by heterogeneous nucleation of transition lattices with a lower energy than the cut zones, then as these overaging structures grow, and the matrix is depleted of solute, the G.P. zones will dissolve as the local solute concentration is reduced below the immiscibility gap for the G.P. zones producing soft zone-free regions. In the present case, since both 7’ and 0’ are heterogeneously nucleated, the question of reversion or overaging appears to invoIve one of competitive


rates, since if the rate of nucleation of the transition

structure is low with respect to the rate of cycling,

the localized softening will be due to reversion, and

where the rate of nucleation of the transition structure

is high, the localized softening will be due to over-

aging. In either case the structure initially present

and responsible for the strength of the alloy is

destroyed.

The rate of nucleation of the transition structure

(0’ or y’ in Al-Cu and Al-Ag respectively) depends

on several factors. One point of paramount import-

ance is that because of minimization of interfacial

energy, due to good lattice matching, 8’ is nucleated

such that {lOO},//{lOO},,. Heterogeneous nucleation

sites on the (100) plane are comparatively rare, and

usually involve climb of the edge component of a

dislocation of strong screwlike character from the

(ill), the deformation plane, to the (100) plane

required for nucleation of 0’.(31) Before nucleation

can occur, such climb must take place as a function

of vacancy concentration, temperature and time.

The dislocation structures observed, for example in

Fig. 5, indicate that conditions in the deformed

regions did not favor the formation of dislocations

into helices, perhaps because of the oscillating motion

of the dislocations. Without appropriate sites, diffu-

sion of copper atoms to nucleate 8’ does not occur.

On the other hand, the y’ precipitate is relatively

easily nucleated. In this case, the orientation re-

lationship is (001 l},,//{l 1 l},,. Interfacial energy

considerations are less important because of the

nearly identical size of the Al and Ag atoms. Appro-

priate nucleation sites of y’ are now quite common,

consisting of faults on (111) (the deformation plane),

which may be initially present or created by the

splitting of a whole dislocation into partials, a process

described in detail by Hren and Thomas.(2*) Since

for y’, there is no climb step involved as for 8’, the

nucleation depends upon the rate of diffusion of Ag

atoms, aided by the vacancies created by the motion

of the dislocations. Thus the rate of nucleation is

many times faster for y’ than for 8’. For a given

cyclic straining, nucleation of y’ is much more probable

than that of 0’.

These considerations appear to be consistent with

the observations and account for the localized

softening which may occur either because of reversion

of the hardening precipitate or the nucleation of

competitive overaging reactions.

CONCLUSION

1. During the cyclic straining of these alloys in the

solution treated condition, the copper and silver

solute atoms limit the substructure formation and

alter the nature of the loops generated in contrast

to the complex dislocation tangles and cell structure

produced in pure aluminum.

2. Interaction of the dislocations produced by

cyclic straining with the metastable hardening struc-

tures in age hardening aluminum alloys can generate

soft regions within the microstructure where further

deformation is concentrated and the probability of

crack nucleation increased. Two types of reactions

may produce these soft regions:

(a) Reversion in Al4 wt. % Cu by repeated cutting

of the G.P. II zones.

(b) Overaging reactions in M-15 wt. % Ag whereby

the nature of the deformation introduced by

cyclic straining nucleates overaging reactions

(formation of the transition lattice y’ and

transformation to equilibrium phase y) causing

solute depletion of the matrix and resulting in

local dissolution of the G.P. zones.

3. The ease of nucleation of the transition struc-

tures, e.g. y’ or 8’, is the primary consideration as to

whether a given zone-strengthened alloy will revert or

overage as a consequence of cyclic straining.

ACKNOWLEDGMENTS

The authors wish to thank L. R. Sefton for assist-

ance on the electron microscope and R. A. D’Angelo

for cyclic straining and specimen preparation.

1.

2.

3.

4.

5.

6. 7.

8.

9.

10.

:::

13. 14.

15.

16.

17. 18.

19.

REFERENCES

J. C. GROSSKREUTZ and P. WALDOW, Acta Met. 11, 717 (1963). R. L. SEOALL, P. G. PARTRIDGE and P. B. HIRSCH, Phil. Mw. 6. 1493 (19611. R. bT. $VILSO& and P. J. E. FORSYTH, J. Inst. Met. 87, 336 (1959). R. L. SEGALL and P. G. PARTRIDGE, Phil. Mug. 4, 912 (1959). fi. L.‘SEOALL, Electron M~CTO~CO~Y and Strength of Metals, p. 515, Interscience Publishers, New York (1963). R. L. SEQALL and J. M. FINNEY, Acta Met. 11,685 (1963). T. BROOM, J. H. MOLINEAUX and V. N. WHITTAKER, J. Inst. Met. 84. 357 (1955-56). T. BROOM, J. A. MA~ZA and ‘v. N. WHITTAKER, J. Inst. Met. 86, 17 (1957-58). J. M. SILCOCK, T. J. HEAL and H. K. HARDY, J. Inst. Met. 82, 239 (1953-54). R. B. NICHOLSON, G. THOMAS and J. NUTTING, J. Inst. Met. 87. 429 (1958-59). G. LEN&R, R‘es. Aeroimutique 36, 35 (1953). D. L. ROBINSON, private communication. Alcoa Research Laboratories, New Kensington, Pa. G. THOMAS and M. J. WHALEN, Phil. Msg. 4, 511 (1959). P. R. SWANN, Electron Microscopy and Strength of Metals, p. 147, Interscience Publishers, New York (1963). G. W. J. WALDRON, private communication. General Electric Co., London, England. R. B. NICHOLSON and J. NUTTING, Phil. Mag. 3, 531 (1958). A. KELLY and M. E. FINE, Acta Met. 5, 365 (1957). J. G. BYRNE, M. E. FINE and A. KELLY, Phil. Mag. 6, 1119 (1961). P. J. E. FORSYTH, J. Aust. Inst. Met. 8, 52 (1963).


20. A. J. MCEVILY, J. B. CLARK, E. C. UTLEY and W. H. 26. E. J. FREISE, A. KELLY and R. B. NICHOLSON, Acta Met. HERRENSTEIN, !f’rans. Amer. Inst. Min. (Metd.) Engrs. 9, 250 (1961). 227, 1093 (1963). 27. A. GUINIER, 2. Metall. 49, 217 (1952).

21. S. KODA, K. MATSUURA and S. TAKAHASHI, J. Inst. Met. 28. J. A. HREN and G. THOMAS, Trans. Amer. Inst. Min. 91, 229 (1963).

22. L. G. BONAR and A. KELLY, Progress in Materials Science, (Metall.) Engm. 227, 308 (1963).

29. A. M. ADAIR, R. E. HOOK and H. J. GARRETT, ARL 132 10, No. 3, 318 (1963). (1961).

23. R. J. PRICE and A. KELLY, Acta Met. 10, 981 (1962). 24. R. B. NICHOLSON and J. NUTTINQ, Acta Met. 9,332 (1961).

30. C. LAIRD, private communication. Scientific Laboratory,

25. R. BAUR and V. GEROLD, Acta Met. 10, 637 (1962). Ford Motor Co., Dearborn, Michigan.

31. G. THOMAS and M. J. WHELAN, Phil. Mug. 6,1103 (1961).

Interaction of dislocations and structures in cyclically strained aluminum alloys

Documents

Transcript of Interaction of dislocations and structures in cyclically strained aluminum alloys