The influence of aging and thermomechanical treatments on the fatigue properties of an Al-6.5 at.% Zn alloy

Materials Science and Engineering, 32 ( 1 9 7 8 ) 41 - 53 © Elsevier S e q u o i a S.A., L a u s a n n e ..... P r i n t e d in t h e N e t h e r l a n d s

41

The Influence of Aging and Thermomechanical Treatments on the Fatigue Properties of an A1-6.5 at.% Zn Alloy

M.-C. LU* a n d S. W E I S S M A N N

Materials Research Laboratory, College of Engineering, Rutgers University, New Brunswick, New Jersey 08903 (U.S.A.) (Received J u n e 16, 1 9 7 7 )

SUMMARY

1. I N T R O D U C T I O N

The effects produced by conventional aging and thermomechanical treatments (TMT) on the microstructures, and their influence on the fatigue properties of an A1-6.5 at.% Zn alloy were investigated. To assess the influence of the microstructures on the fatigue properties, all alloys were aged to give the same initial static yield stress. Cycling induced in the conventionally aged alloys an excessive deformation structure at grain boundaries, which was associated with the formation of precipitation-free zones (PFZ's), and an inhomogeneous, transgranular, banded dislocation structure, which resulted from the cutting o f G.P. zones by moving dislocations. By contrast, cycling induced in TMT alloys a deformation structure uniformly distributed throughout the grains. These alloys exhibited for high-cycle fatigue a tenfold increase in fatigue lives and a 15% increase o f the endurance limit. The improvement of the fatigue properties by TMT was attributed to an effective dispersal o f slip, thereby increasing the resistance to crack nucleation. Fatigueinduced boundary migration was found in specimens aged to contain predominantly G.P. zones when subjected to high-cycle fatigue. The driving force was attributed to the strain energy o f dislocations, which accumulated excessively in PFZ's adjacent to the grain boundaries.

It is generally recognized that the fatigue strengths of precipitation-hardened aluminum alloys are unusually low when compared with their tensile strengths [1 - 3]. This phenomenon has been ascribed to an instability of the metastable metallurgical structure, and has been attributed either to overaging of the precipitates [1 - 4], or to re-solution of the metastable precipitates [4 - 8 ] , or to heterogeneous aging in the form of precipitate-free zones (PFZ's) at grain boundaries and throughout the matrix [9 - 11]. The past several years have witnessed an increased interest in thermomechanical treatment (TMT) to optimize the mechanical and also the corrosion-related properties of precipitation-hardened aluminum alloys [12 15]. The aim of TMT is to produce a metallurgical, stable structure that supports homogeneous deformation, and to achieve this end even multi-staged heat treatments combined with mechanical deformation have been used. Applying TMT to aluminum 7075, Ostermann [16] reported an increase of 25% in the 107 cycles-to-failure stress level when compared with the conventionally aged alloy of the T651 condition. The TMT consisted of room-temperature deformation of the partially-aged alloy and post deformation aging. Reimann and Brisbane [17], studying notched 7075 aluminum, found that the fatigue-life curves remained unchanged by TMT. These authors suggested that TMT may retard crack initiation by strengthening the matrix through slip dispersal b u t may have no effect on crack growth. Recently, however, Mehrpay e t al. [18] showed that in the lowstress region of 7075 aluminum, TMT can

*This p a p e r is b a s e d o n a thesis s u b m i t t e d b y M.-C. L u in partial f u l f i l l m e n t o f t h e r e q u i r e m e n t s for a Ph.D. degree o f R u t g e r s University.

42

promote increased resistance to both crack initiation and crack propagation. Although it is generally agreed that the beneficial effects of TMT derive from a strengthening o f the microstructure, impeding the motion of dislocations, the details of this strengthening effect are by no means clearly understood. Thus, while Ostermann [16] attributes the improved fatigue properties to the heavy entanglement and to the high density o f dislocations induced by cold work, Avitzur [19], in a discussion of Ostermann's work, points out that the major contribution of TMT to improving tensile ductility, fatigue, and stress-corrosion lives, may well be due to the elimination o f the weak PFZ's. The aim of this study is to clarify the influence of TMT on the microstructural strengthening of precipitation-hardened alloys. An aluminum-6.5 at.% Zn alloy was selected as a study material, since preceding studies were able to define the lattice defects introduced in this alloy by conventional heat treatments [20], and also were able to show how these pre-existing lattice defects affected the defect structure generated by cyclic straining [21]. To establish a meaningful basis for comparing the effects of the different microstructures formed by either conventional aging or TMT on the fatigue properties, it was decided that all alloys, regardless of the aging treatment, were to be prepared in such a way as to exhibit the identical static yield stress prior to cycling. This binding stipulation aimed to elucidate the nature and distribution of that initial precipitate structure which would be most efficacious in producing the desired improvements of the fatigue properties. Thus, the different, initial strength levels that different heat and aging treatments usually produce do not have to be taken into account when the effectiveness of the microstructures

on fatigue properties is being compared, since all specimens would start with the same static yield stress.

2. EXPERIMENTAL

Cylindrical specimens with a diameter o f 3.8 m m and a gauge length of 10 mm were prepared from hot-rolled, 1.27 cm thick sheets of an A1-6.5 at.% Zn alloy obtained by courtesy of the Reynold Metal Company, Richmond, Virginia. The composition in weight percent., including trace elements, is given in Table 1. All specimens were solutionized at 550 °C for 2 h in a quartz tube under argon atmosphere and quenched in water. After quenching, the specimens were kept at room temperature for at least 1 day. T h e y were then subjected to either conventional aging (specimens A and B) or to TMT (specimens C and D) to produce different microstructures. The aging parameters, temperature and time, as well as the parameters of mechanical deformation for the TMT specimens, precycling and tensile deformation, respectively, were so selected as to result in nearly identical values o f the static proof stress a0.: = 13.4 kg mm -2. Table 2 lists the aging and TMT treatments of the alloys after the solution treatment. All fatigue tests and precycling treatments were run at room temperature on a TatnallKrouse push-pull axial loading test machine with constant stress amplitude and a frequency of 20 cycles s-:. To characterize the microstructure of the alloys before and after cycling, foils for transmission electron microscopy (TEM) were prepared. They were obtained by slicing thin discs from the bulk specimens perpendicular to the tensile axis, reducing their thickness to about 250 pm with the aid o f a spark-cutting machine,

TABLE 1 Composition of A1-6.5 at.% Zn in wt.% Elements

Al

Zn

Si

Cu

Ti

Fe

Mn, Mg, Ni, Cr

Wt.%

85.41

14.39

0.05

0.05

0.01

0.05

nil

43 TABLE 2 Aging and TMT of alloys after solution heat treatment Specimen designation

Type of aging

Deformation prior to aging

Aging sequence

A B C

Conventional Conventional TMT

150 °C/100 h + 100 °C/24 h 150 °C/20 min + 100 °C/18 h 150 °C/20 min + 100 °C/18 h

D

TMT

None None 100 000 cycles with stress amplitude of 1/3 of elastic limit Tensile strained 4.5% at 150 °C

followed by jet-electropolishing. An electrolyte solution of 8% nitric acid and 92% methanol was used at --30 °C. In examining the foils, the electron microscope was operated at 200 kV.

3. RESULTS

To demonstrate the influence of the diverse microstructures on the fatigue properties, the scheme of presenting the results will adhere to the following sequence: (i) characterization of microstructures produced by conventional heat treatment, (ii) microstructures induced by TMT, (iii) effect of microstructures on fatigue properties, (iv) characterization of microstructures after cycling. (i) Characterization o f microstructures produced by conventional heat treatment (a) Duplex structure -- Specimen A Previous studies have shown that the defect structure induced in this alloy after solution treatment consists of quenched-in dislocation loops, helical dislocations and reaction products of these defects [20, 21], as well as of G.P. zones [22, 23]. Experiments have shown that maximum strength and hardness are reached after short aging at room temperature, and that the strength increases are principally due to the coherency strains generated by the fully developed spherical G.P. zones [23]. When the alloy was aged at higher temperatures (but below 140 °C, which corresponds to the solvus temperature of the G.P. zones, viz., 100 or 123 °C), the size of the G.P. zones became larger and their shape changed from spherical to ellipsoidal [24]. Simultaneously, a heterogeneous precipitate, incoherent with the matrix, emerged.

100 °C/15 h

When the specimen was aged at 150 °C, which is a little above the G.P. solvus temperature, the G.P. zones dissolved completely in 5 min, while the secondary defects annealed completely out in 1 h [25]. After 1 h aging at 150 °C, the formation of plate-like, incoherent precipitates began to occur with orientation parallel to (111) of the matrix. The density of the precipitates increased gradually with increasing aging time and reached a maximum after about 20 h aging. The size also continued to increase. After 100 h aging, the morphology of the precipitates changed from plate-like to parallelepiped shape. Finally, the precipitates changed their shape from a parallelepiped to a prismatic appearance. Figure 1 shows the isothermal sequence of morphology resulting from aging at 150 °C for various aging times. Since 150 °C is above, and 100 °C below, the G.P. solvus temperature, two types of precipitates were obtained when the specimen was aged at 150 °C/100 h + 100 °C/24 h. One type of precipitate was of small size, about 80 A in diameter, and consisted of G.P. zones, and they were most likely homogeneously nucleated. The other type of precipitate was plate-like in shape. They were about 800 to 1 000 h in length and were characterized by an h.c.p, structure. They nucleated heterogeneously at the sites of quenched-in loops and dislocations, and their preferred growth was parallel to the (111) of the matrix. The density of large precipitates was found to be of the order of the density of quenched-in loops, ~1015 c m - 3 . The electron diffraction pattern of Fig. 2(a) and the schematic illustration of Fig. 2(b) disclose the h.c.p, structure of the large-size precipitates as well as the orientation relationship, which was (0001)ppt//(lll) matnx. Based on the mea-

44

D

(a)

0"

"-

~2 0,~200110 •

. MATRIX e. PPT

(b) Fig. 1. Microstructures of precipitates resulting from aging at 150 °C. (a) 20 h; (b) 100 h; (c) 320 h.

surements of the interatomic spacings of the diffraction pattern, it was deduced that the composition of the plate-like precipitates was close to that of pure zinc. It was not possible, however, to demonstrate the existence of the G.P. zones by electron diffraction. A typical distribution of the precipitates in this alloy, aged to the duplex structure, is shown in Fig. 3. One may note that the platelike precipitates were associated with quenched-in dislocation loops, homogeneously distributed. These loops resulted from the aggregation of vacancies the concentration of which became supersaturated when the alloy was quenched from the high temperature of the solution treatment [20, 21]. In regions of the specimens which contained a high dislocation density, presumably induced by the quenching stresses, the morphology of the plate-like precipitates differed from the general morphology. Figure 4 may serve to illustrate the topography of

Fig. 2. Diffraction pattern corresponding to the microstructure of Fig. l(b). Z.A. of matrix [111]; Z.A. of precipitate [0001 ]. (a) Electron diffraction pattern; (b) schematic illustration.

Fig. 3. Distribution of precipitates in heat-treated alloy A associated with quenched loops. Z.A. [111 ].

precipitates with regard to the stress-induced dislocation structure. Although the size of precipitates remained essentially the same, the precipitates segregated preferentially on dislocation sites, while regions adjacent to the

45 vacancies, greatly reduced the density of quenched-in loops [21]. Consequently, these sites were not available for nucleation of the precipitates. As shown in Fig. 5, precipitation was also found at grain boundaries, but the zone adjacent to the grain boundary was characterized by the absence of the large type of precipitates. This zone will be referred to as the precipitation-free zone, or PFZ. (b) Microstructure w i t h G.P. z o n e s -Specimen B

Fig. 4. Distribution of precipitates in heat-treated alloy A for volume containing a high density of dislocations induced by quenching strains.

Fig. 5. PFZ in heat-treated alloy A.

In order to obtain a microstructure with a very low volume fraction of the zinc-rich precipitate and a G.P. zone formation that would induce the same initial yield stress as the previously aged specimen, the following heat treatment was carried out. The specimen was first heated at 150 °C for 20 min to dissolve completely the initial G.P. zones which may have formed non-uniformly. Subsequently, the specimen was re-aged at 100 °C for 18 h. The G.P. zones reappeared at 100 °C, because this temperature was below the G.P. solvus temperature but high enough to supply the necessary diffusion mobility for nucleation of G.P. zones. Therefore, the specimen with 150 °C/20 min + 100 °C/18 h aging treatment contained mainly G.P. zones, presumably more uniformly distributed. Generally there were no observable PFZ's close to the grain boundaries. However, large dislocation loops originating from the quenched-in prismatic loops were clearly discernible. The density of these loops was low, and near the grain boundary they were entirely absent, as may be seen from Fig. 6.

(ii) Microstructures i n d u c e d by T M T -S p e c i m e n s C and D

Fig. 6. Region at grain boundary with PFZ and prismatic dislocation loops in heat-treated alloy B. Z.A. [123].

dislocations were characterized by a low density of precipitation. The reason for this phen o m e n o n is attributed to the fact that the dislocations, acting as a strong sink for

Two kinds of TMT were performed in this study. They consisted of: (1) precycling at room temperature and (2) tensile deformation at elevated temperature. Both treatments were followed by aging. In the TMT pertaining to precycling, the quenched specimen was cycled for 100 000 cycles with a low stress amplitude corresponding to about one-third of the yield stress of the quenched specimen. In order to provide a comparable a m o u n t o f G.P. zones as in the conventionally aged specimens, the specimen

46

was subsequently heated at 150 °C for 20 min to dissolve the G.P. zones formed from quenching, and was again aged at 100 °C for 18 h to re-form the G.P. zones, thus completing the TMT cycle. TEM investigations showed that the precycling had a marked effect on the dislocation structure in the vicinity of grain boundaries, giving rise in these areas to a dense accumulation o f dislocations as shown in Fig. 7. Since these dislocations appear to be pinned by a fine precipitate structure, the TMT with precycling was expected to result in a strengthening effect of the grain boundaries. Although the dislocation structure of the grain interior was n o t substantially altered by the precycling operation, an important strengthening effect was nevertheless expected from the excess vacancies that were generated by the low stress cycling [26] and which should enhance the nucleation of homogeneously distributed precipitates. Thus, this type of TMT was designed to strengthen all those regions of the specimen that would appear " s o f t " by conventional heat treatment. In the second kind of TMT, the quenched specimen was first strained 4.5% in tension at 150 °C followed by aging at 100 °C for 15 h. It will be noted that the later aging process was similar to the post-aging operation of the TMT containing the precycling step, except for the somewhat shorter aging time. This adjustment of aging time became necessary so as to insure the same initial static yield stress for all the various microstructures investigated. The deformation temperature requires special consideration in TMT. If the specimen was deformed at a temperature below the G.P. solvus temperature, for example, at room temperature, the deformation became highly localized, resulting in an inhomogeneous distribution of dislocations characterized by widely spaced slip bands. By contrast, tensile deformation above the G.P. solvus temperature resulted in a homogeneous distribution of dislocations throughout the specimen, including regions adjacent to the grain boundaries, as shown in Fig. 8. Since the specimen was subjected to a macroscopic, plastic strain, the dislocations and vacancies were effectively introduced in all grains of the specimen. The uniform dislocation structure was stabilized by the post-aging step. In this way, the process of nucleation of precipita-

Fig. 7. Dislocation structure at GB in TMT alloy C

induced by precycling, replacing the " s o f t " PFZ zone. Z.A. [211].

Fig. 8. Uniform dislocation structure pinned by fine precipitate in TMT alloy D.

47

tion in denuded zones was accelerated and, consequently, soft regions at grain boundaries and in the grain interior were removed. Thus, the specimens treated by this type of TMT contained not only a uniform, fine-precipitate structure similar to that obtained by the precycling TMT, but also a uniform, stabilized dislocation structure.

(iii) Effect of microstructures on fatigue properties After the various, developed microstructures were characterized by TEM and tested for equivalent static tensile properties, they were subjected to dynamic cycling. The results of the fatigue tests are shown in the plot of Fig. 9, giving the relation between applied stress and number of fatigue cycles to fracture (S-N curves). It will be seen that the two specimens subjected to TMT {specimens C and D o f Table 2) exhibited superior fatigue properties to those aged by conventional heat treatment (specimens A and B).

,. E

9

~8 E

o

+

o

•

*

•

i' 6

10 4

10 s

10 ~

10 7

Cycles t o f a i l u r e

Fig. 9. Dependence of fatigue life on aging and TMT. Initial yield stress for all specimens 13.4 kg mm -2 + = S p e c i m e n A ; o = B ; ~ = C ; • =D.

The improvement of fatigue properties by TMT is particularly noticeable for high-cycle fatigue life when the applied stress was small, and is less conspicuous for low-cycle life when the applied stress was large. Thus, it may be seen that for the applied stress o f -+ 7.2 kg mm -2, the fatigue life of the TMT specimen was extended by a factor of 10. Expressed in terms of the effect on the endurance limit, the TMT specimens exhibited an increase of about 15% over the conventionally heattreated specimens A and B. With regard to the latter, specimen B, containing principally G.P. zones, displayed the

poorest fatigue properties. Compared with specimen A, the deleterious effect on the fatigue properties of specimen B was particularly noticeable in high-cyclic stress fatigue, while for low-cycle stress fatigue the difference between specimen A and specimen B gradually disappeared.

(iv) Characterization of microstructures after cycling To establish the effect of microstructure on the fatigue properties, it was necessary to characterize the structures after dynamic cycling and to compare them with those prior to cycling and also, in certain instances, with those subjected to unidirectional deformation.

(a) Conventionally aged specimens In alloys A, aged for the duplex structure of zinc particles and G.P. zones, cyclic deformation induced inhomogeneous deformation if a stress amplitude below the macroscopic yield point was employed. The inhomogeneity o f deformation varied, not only from grain to grain, but also within the individual grains. Indeed, some grains showed virtually no deformation at all, while others were strongly affected by it. This variation was crystallographically conditioned, and was due to the unequal resolved shear stress applied to the slip systems of the individual grains. When the specimens were fatigued with the stress amplitude o ~> + 9.6 kg mm -2, which corresponded to 0.72 of the macroscopic yield stress o0.1, a high concentration of dislocations was generated in those regions which pertained to the PFZ's in the uncycled specimens, as may be seen by comparing Fig. 10 with Fig. 5. With this applied stress, the dislocations which were principally generated at or near the grain boundaries were also able to penetrate beyond the regions formerly occupied by the PFZ's and to distribute themselves uniformly in the interior of the grain. There was no evidence that either the size or distribution of the zinc particles were affected by the movement of the dislocations during the fatigue process, and it must be concluded, therefore, that the zinc particles controlled the bypass mechanism of the moving dislocations. It was observed that concomitant with the preferred accumulation o f dislocations in the vicinity of the grain boundaries, there oc-

48

Fig. 10. Fatigue-induced dislocation structure o f alloy A. Cyclic stress amplitude a = -+ 9.6 kg m m - 2 ; N = 42 300 cycles.

curred an accelerated growth of the zinc particles. This accelerated growth appeared to be facilitated b y the dense dislocation entanglements at the grain boundaries. When the specimen was cycled with a low stress amplitude, o < + 7.2 kg mm -2 (below 0.54 of yield stress Oo.1), most of the area inside the grain remained unaffected by the fatigue process. As in cycling with the higher stress amplitude, the dislocations were also generated at, or near, the grain boundaries, b u t in contradistinction to the high stress cycling they were unable to proceed b e y o n d the PFZ. They entangled and remained confined either at sites of the PFZ's or at the border between the PFZ and the grain interior containing the dispersed zinc particles. This t y p e of interaction resulted in a banded dislocation structure, and such bands appear to have originated either from dislocation sites at the grain boundaries or from grain triple points, in agreement with the observations made b y Kiritani and Weissmann [21]. Accumulation of dislocations at grain boundaries, and the formation o f a banded dislocation structure (Fig. 11), became even more pronounced in B specimens, consisting principally of G.P. zones, when cycled with a stress amplitude o = + 9.6 kg mm -2. A remarkable deformation response pertaining to the low stress cycling of these alloys concerns the observation of fatigueinduced grain boundary migration. It was observed that these alloys exhibited, prior to cycling, grain boundaries which were straight or smoothly curved. After cycling,

Fig. 11. Banded, transcrystalline dislocation structure induced in cycled alloy B, o = -+ 9.6 kg m m - 2 . Z.A.

[211].

Fig. 12. Grain b o u n d a r y migration (shown by arrow) induced in cycled alloy B, a = t 7.2 kg m m - 2 . IB = Initial b o u n d a r y position; P = precipitate particles f o r m e d o n the b o u n d a r y during migration.

the boundaries frequently showed an irregular, sinusoid shape such as that shown in Fig. 12. This boundary shape seemed to have arisen from an accommodation process of the t w o

49 moderately strained, neighboring grains involving boundary migration. The boundary with the initial position marked IB had migrated after cycling to the new sinusoid position where large, zinc-rich particles were precipitated, designated in Fig. 12 by the letter P. The distance traversed by the boundary was of the order of 1 gm and appears to correspond to the width of the PFZ. Both the fatigue-induced boundary migration and the enhanced precipitation on the migrated boundary have been previously observed on the same alloy aged at room temperature [21]. The grain boundary migration was characteristic of low-stress cycling. It was not observed at high-stress cycling nor was it observed in the alloys aged for duplex microstructure.

Fig. 13. Microstructure of cycled TMT specimen C, a = ± 7.2 kg mm-2. Z.A. [110].

(b) TMT specimens There exist some striking differences in the fatigue-induced microstructures among specimens which were conventionally aged and those which were aged by TMT. It has been shown that for conventionally aged specimens, cycling gave rise to a preferential deformation o f regions associated with grain boundaries. When a low stress amplitude was employed, specimens aged by TMT did not exhibit such preferential deformation sites. Thus, when a cyclic stress, o ~< + 7.2 kg m m - 2, was applied to TMT specimens, the fatigueinduced deformation was homogeneous also for regions adjacent to the grain boundaries, as shown in Fig. 13, and bands of piled-up dislocations, deleterious to the mechanical

properties, did not form. However, inhomogeneous deformation characteristics inside the g a i n s were still observed in TMT specimens when a high stress amplitude was applied, giving rise to localized accumulations of dislocations in the grain interior. These were formed, presumably, by the cutting of G.P. zones. It was also observed that the local sites of dislocation accumulation inside the grains were less frequent and less pronounced in the tensile-strained than in the precycled TMT specimens. Grain boundary migration was never observed for TMT specimens. Table 3 may serve to summarize the fatigue-induced deformation characteristics and their dependence on stress level.

TABLE 3 Dependence of fatigue-induced deformation characteristics on stress level Specimen

Heat treatment

Stress level o> -+ 9.6kgmm -2

o<-+ 7.2kgmm -2

A

150 °C/100 h + 100 °C/24 h Duplex structure

Inhomogeneous distribution o f dislocations at grain boundaries

Preferred deformation at PFZ's

B

150 °C/20 min + 100 °C/18 h Principally G.P. zones

Dislocation bands in grain interior and dislocation accumulation at grain boundaries

Preferred deformation

C

TMT -- 100 000 precycles + 150 °C/20 rain + 100 °C/18 h

Localized accumulation of dislocations in grain interior

Homogeneous distribution o f dislocations

D

TMT -- tensile-strained 4.5% at 150 °C + 100 °C/15 h

Some localized accumulation o f

dislocation in grain interior

Homogeneous distribution o f dislocations

at PFZ's. Grain boundary migration

50

(c) Fracture surface studies Fracture surface studies by scanning electron microscopy (SEM) were carried o u t with the aim of obtaining some inferences concerning the relative amount of crack nucleation and crack propagation in the fatigue life of the specimens. In particular, it was desirable to ascertain to what extent TMT would affect either one of these processes. The scanning electron micrographs obtained had some general features in common. One could distinguish rough regions in which the fracture surfaces delineated sharply the contours of the grain boundaries. These intercrystalline markings were interpreted as topographical features associated with the last stage of crack propagation when the specimens began to fail catastrophically as a result of overloading. The regions of greatest interest, however, were those which exhibited transcrystalline fatigue striations such as those shown in Fig. 14. Following the interpretation of Broek [27, 28], the spacings between

Fig. 14. SEM t r a n s g r a n u l a r fatigue s t r i a t i o n s o n fract u r e surface o f cycled s p e c i m e n B. Cyclic stress amplit u d e o = + 7.2 kg m m - 2 ; N = 62 0 0 0 cycles.

~hese fine striations were taken to be indicative of the average distance traversed by the crack per fatigue cycle. The fractional areas which exhibited these fine striations were comparable in all specimens and were 25 30% for the conventionally aged ones and slightly less for those subjected to TMT. The measurements o f the average striation spacings yielded an average propagation rate of ~ 5 × 10 -4 mm/cycle for all specimens. On the basis of these measurements, an order-of-

magnitude estimate could be made regarding the relative contribution of the crack propagation to the fatigue life of the specimens. For the application of a fatigue stress, o = + 7.2 kg mm -2, the crack propagation amounted to about 15% or less o f the life for virtually all the specimens investigated.

4. DISCUSSION

(i) Effect o f precipitate distribution on slip dispersion and fatigue life It has been shown that when the alloys were subjected to a conventional heat treatment, a matrix was produced which was agehardened inhomogeneously. PFZ's close to the grain boundaries as well as in the interior were formed which, u p o n cycling, became preferred sites for plastic deformation (Figs. 5, 10). The cause of the transgranular deformation bands (Fig. 11) was the G.P. zones, which represent weak obstacles to the moving dislocations. They were easily cut and finally destroyed. The strength o f the alloy, which depended on the presence of the coherent G.P. zones, was thus locally reduced. These sheared G.P. zones facilitated the subsequent passage o f moving dislocations, which resulted, thus, in slip confined to narrow, transcrystalline deformation bands. Such deformation bands are responsible for excessive stress concentration at grain boundaries and favor crack nucleation. By contrast, the specimens aged to produce a duplex precipitate structure (specimen A) exhibited, after cycling with a large strain amplitude, a dislocation distribution which was more homogeneous in the grain interior than that of alloy B. The cause of the more uniformly distributed slip was the incoherent zinc particles which were uniformly distributed inside the grain and acted as effective obstacles to the moving dislocations. The characteristics of homogeneous slip in the microstructure explain the better fatigue life for this specimen at relatively high stress fatigue (about twice the life of specimen B, Fig. 9). However, when the fatigue stress was low, i.e., o <<.+ 7.2 kg mm -2, the structural inhomogeneities o f the soft PFZ's became the dominant factor in the production o f inhomogeneous slip, since fatigue deformation at low stresses involves only the move-

51 ment of dislocations in a "to-and-fro" manner over short lattice distances within the grain. The dislocations generated from grain boundaries were unable to proceed b e y o n d the PFZ's due to the resistance of zinc particles located at the boundaries of the PFZ's. As a result of this development, the regions adjacent to the grain boundaries, comprising the PFZ's, became deformed preferentially, while the rest of the matrix remained virtually unaffected. This inhomogeneity in plastic deformation would favor crack nucleation, and was probably principally responsible for the poor performance of this aged alloy in high-cycle fatigue. As may be seen from Fig. 9, the fatigue life for cycling with low stress amplitude led practically to the same value as that obtained for the alloy aged only for G.P. strengthening. By contrast with the conventionally aged specimens, the TMT specimens exhibited a marked increase in fatigue life when cycled with a low stress amplitude (Fig. 9). The precycled as well as the tensile-prestrained TMT specimens displayed, after cycling with a low stress amplitude, a homogeneously uniform dislocation structure throughout the grain, and the banded microstructure, observed in the conventionally aged specimens, was totally absent. These results are in agreement with those obtained b y Stubington [29], who showed that a finer and more uniformly distributed precipitate was obtained if aging followed cyclic loading of a supersaturated solid solution o f an aluminum alloy. They were also in agreement with the results o f McEvily e t al. [8], who found that the fatigue life o f a 2024-T4 aluminum alloy was markedly increased if the specimen was re-aged at 150 °C for 16 h after cycling for 10% o f its fatigue life at low stress. The authors attributed this beneficial effect to the re-aging o f coherent precipitates which, as a result of cycling, underwent a reversion process. The beneficial results o f the TMT obtained in this study were interpreted, however, in terms of entirely different concepts from those advanced by McEvily e t al. [8]. The precycling step as well as the unidirectional deformation step were purposely interposed in TMT to insure the introduction of dislocations homogeneously distributed throughout the specimen. These dislocations were conceived to function as sites at which preferred

precipitation o f the particles would take place upon subsequent aging. The obstacle effect of the precipitate particles for the moving dislocations generated by cycling depends, besides their inherent strength and shear modulus, on the proper interparticle distance. The spacings of the interparticle distance can be so designed in the microstructure that it is energetically more favorable for the dislocations to bypass the particles than to shear them [30]. If the interparticle distance is too short the particles will be cut by the dislocations, resulting in inhomogeneous slip and concomitant dislocation pileup. If, on the other hand, the interparticle distance is t o o large, slip dispersal may take place; but the obstacle effect will be reduced and, hence, the effective strength of the precipitate structure will be diminished. Consequently, the dislocation structure introduced in TMT by precycling or unidirectional extension prior to aging serves to set the blueprint for the particle distribution that is to be developed. By varying the number of precycles or strain of tensile deformation in the TMT procedure, the future interparticle distance can be controlled. Moreover, since upon aging the zinc-rich particles will precipitate preferentially on dislocation tangles, they become anchored and thereby become more effective obstacles when compared with particles precipitated in the absence of an underlying dislocation structure, such as in conventional aging. It is interesting to note that the unidirectional tensile deformation o f 4.5% at 150 °C in TMT introduced a dislocation structure which was locally more heavily tangled than that produced by precycling. Yet, the effect of both TMT's on fatigue life and fatigue strength was quite equivalent. At high stresses, inhomogeneous slip bands with pronounced accumulation of dislocations were still observed in both types of TMT. These results suggest that it was not the dislocations p e r se which were responsible for the improvement o f the fatigue properties, b u t that the strength increases stemmed from the particles and their distribution [19]. Therefore, the dislocation structure introduced by TMT should be considered, primarily, as a means of facilitating the homogeneous distribution of particles and of controlling the interparticle distance.

52

One of the most effective contributions made by TMT toward the extension of the high-cycle fatigue life was the removal of structural inhomogeneities such as the PFZ's at grain boundaries and within the matrix. Since such soft regions were absent due to the controlled dispersal of slip, preferential deformation, which otherwise would have taken place in such soft regions and would have led to crack nucleation, was avoided. Since the fracture surface analysis of the fine, fatigue-induced striation bands did not show any marked difference between the TMT and conventionally heat-treated alloys, it was concluded that for the laboratorymade alloys of this study, the beneficial effect o f extending the life in high-cycle fatigue by TMT could not be attributed to an increase of resistance to crack propagation b u t must be ascribed, rather, to an increase of resistance to crack nucleation. In commercial alloys, however, such as 7075-A1, which contain inherently stronger precipitate particles as well as dispersed, second-phase particles, the TMT may very well be capable of generating a microstructure which could also extend the crack propagation stage [18].

(ii) Fatigue-induced grain boundary migration Fatigue-induced grain boundary migration was found only at low stress fatigue in specimens conventionally aged for G.P. zones. It appears that the driving force was provided by the strain energy of the dislocations which, generated by cycling, accumulated preferentially in the PFZ's or denuded zones adjacent to the grain boundaries. In addition, the excess concentration of vacancies, generated by low cyclic stress, may have enhanced, also, the diffusion rate necessary for grain boundary migration [26]. The extent of the grain boundary migration was limited to the width of the PFZ's (Fig. 12), in agreement with previous studies by Kiritani and Weissmann [21]. Several factors appear to be responsible for this limited distance in boundary migration: (1) the presence of G.P. zones and incipient fine precipitates of zinc particles located at the boundaries between PFZ's and the grain interior, (2) retardation caused by enhanced precipitation of zinc particles at grain boundaries,

(3) increase of surface energy accompanying the migration process as manifested b y the sinusoidal shape of the grain boundaries after migration (Fig. 12).

5. C O N C L U S I O N S

The effects produced by various aging and thermomechanical treatments (TMT) on microstructure and fatigue properties of polycrystalline A1-6.5 at.% Zn alloy were investigated. To assess the influence of the different microstructures resulting from the diverse treatments on the fatigue properties, the treatments were so designed as to produce the identical static yield stress prior to cycling, o0.1 = 13.4 kg mm -2. The following results were obtained: (1) The nature and extent of the deformation induced by cycling depended on the aging treatment and on the cyclic stress amplitude applied. For a comparable aging treatment and stress amplitude, the alloys conventionally aged became more deformed at grain boundaries. The microstructural evidence showed that dislocations were generated from sources at grain boundaries, and that matrix or transgranular deformation originated from grain boundary regions. (2) The excessive deformation of soft regions at grain boundaries in specimens subjected to conventional aging was associated with the formation of precipitation-free zones (PFZ's) resulting from the aging treatment. (3) A non-uniform, transgranular, banded dislocation structure was the typical deformation feature of the microstructure o f cycled specimens aged to contain G.P. zones without zinc particles. These structural inhomogeneities derived from the cutting of the G.P. zones by the moving dislocation during the cycling process. (4) Fatigue-induced grain boundary migration was found only in specimens aged to contain predominantly G.P. zones when a low- or medium stress amplitude was employed. The driving force was attributed to the strain energy o f the dislocations, which accumulated excessively in the PFZ's adjacent to the grain boundaries. (5) Specimens subjected to TMT exhibited a tenfold increase in fatigue life and a 15% increase of the endurance limit at the level of

53 106 c y c l e s w h e n a c y c l i c s t r e s s o f o = + 7 . 2 k g mm -2 was employed. (6} T h e i m p r o v e m e n t o f t h e f a t i g u e p r o p e r ties by TMT was attributed to a more effect i v e d i s p e r s a l o f slip. T h i s , in t u r n , w a s a c c o m plished by a more homogeneous distribution of precipitated zinc particles and an effective interparticle distance, both of which were conditioned by the dislocation structure i n t r o d u c e d p r i o r t o aging. (7) Fracture surface studies of fatigue striations supported the interpretation that the increase of fatigue life by TMT in the laboratory-made alloys of this study was primarily due to an increased resistance of the alloys to crack nucleation.

ACKNOWLEDGEMENTS

The support of this work by the Rutgers Research Council is gratefully acknowledged.

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