!!Effects of alloy modification and thermomechanical processing on recrystallization of al mg-mn...

Effects of Alloy Modification and Thermomechanical Processing on Recrystallization of AI-Mg-Mn Alloys

K. KANNAN, J.S. VETRANO, and C.H. HAMILTON

The 5083 A1 alloy (Al-4.75Mg-0.8Mn) holds potential for superplastic forming (SPF), but slow rates of forming limit its use for many applications. Higher strain rates are believed possible through the development of finer grained microstructures or stabilized subgrain structures. Grain sizes after recrystallization and recrystallization characteristics are known to be dependent on the amount and distribution of second-phase particles in the matrix. In this study, the concentration and sizes of such particles were varied by additions of particle-forming elements of Mn and Zr and by modifications of the rolling and aging schedules (thermomechanical processing (TMP)). The investigation involved studying recrystallization kinetics at different temperatures and correlating the grain sizes with particle sizes and volume fractions. The addition of Mn and Zr, for the composition ranges and TMP methods studied, resulted in a substantial reduction of the recrystallization kinetics, but complete suppression of static recrystallization (or subgrain stabilization) was not observed. However, statically recrystallized grain sizes as small as 6 txm were achieved.

I. INTRODUCTION

THE A1-Mg-Mn alloy 5083 is of interest for superplastic forming (SPF) because of its moderate strength, corrosion resistance, weldability, and because it has been shown to be sufficiently superplastic to permit SPF of components, t~- 61 Elongations of over 300 pctt2.3] have been achieved using commercial grades of the 5083 alloy. Development of SPF- grade 5083 alloys, chiefly by reducing the Fe and Si im- purity content, has led to higher elongations of up to 630 pct. H,51 Other modifications have involved additions of Cu, Zr, and Mn.t4.61 Vetrano et al.t3] also studied the effect of thermomechanical processing (TMP), aimed at reducing grain sizes, on superplastic elongation in commercial alloys. This study, which focused on the effect of cold rolling reductions, concluded that thickness reductions of at least 5:1 gave the best superplastic tensile elongations.

In most cases, the strain rates necessary to obtain ac- ceptable ductility are in the range of 2 • 10 4 S-I to 4 • 10 -4 S - I , requiring relatively long forming times, of the order of 30 minutes to a few hours. There is, therefore, interest in increasing the superplastic strain rate of this type of alloy through further alloy development and TMP, while retaining the low-cost casting methods for the material. There has been some success in achieving elongations of 500 to 600 pct at high strain rates of 10 -2 s -1 in the A1- 5Mg-0.8Zr alloy system processed by powder metallurgy (PM). t71 It is easier to obtain a finer dispersion of second- phase particles and finer grain sizes by PM than by ingot casting, but the former is substantially more expensive.

It is well known that a primary factor in producing superplasticity is a fine grain size, and the finer the micros-

K. KANNAN, Graduate Student, and C.H. HAMILTON, Professor, are with the School of Mechanical and Materials Engineering, Washington State University, Pullman, WA 99164-2920. J.S. VETRANO, Senior Research Scientist, is with the Structural Materials Research Section, Pacific Northwest National Laboratory, Rlchland, WA 99352.

Manuscript submitted October 30, 1995.

tructure, the higher the strain rate for superplastic deformation (SPD). Two types of microstructures are believed to be suitable for improved SD rates: (1) stabilized subgrain structure leading to continuous dynamic recrystallization during SPD and (2) statically recrystallized equiaxed microstructure with grain sizes less than about 5 /xm, which is stable at temperatures greater than 0.5Tin (in degrees kelvin). The former has been demonstrated in the 8xxx A1-Li-Zd 8'91 and the Supral 2xxx t~~ series of alloys (by ingot casting), but there have not been many such studies with the 5083 series. Almost all studies reported for this alloy have been with (2), i.e., statically recrystallized microstructures, with grain sizes of about 8 to 10/zm.

Recrystallization characteristics and resulting grain sizes in A1 alloys, arising out of a particular TMP, are related to the amount and distribution of second-phase particles in the matrix. Hard, nondeformable particles, larger than a critical size, can serve as sites for recrystallization and, thus, influence the as-recrystallized grain size Dro c. For most AI alloys, this critical particle size has been observed to be around 1 to 2 /xm.[~ ~,~21 An increase in the density of such nucleation sites leads to finer recrystallized grains. How- ever, another aspect of the effect of particles, viz. the Zener drag pressure, can sometimes lead to a contrary effect on recrystallization and growth, t13,14̀~sl Particles of all sizes exert a Zener drag pressure P~ on moving boundaries, restricting their migration, where

t 9_ = t9 y f / r [1]

where y is the surface energy,/3 is a constant, a nd fand r represent the volume fraction and size of the particle.

Both subgrain boundaries in the unrecrystallized sub- structure and grain boundaries of new recrystallized grains are subject to Zener drag, resulting in recrystallization and grain growth being influenced by particle dispersion. This leads to opposing effects on the final grain size. Drag pressure on subgrain boundaries by the fine particles leads to a coarser D .... by preventing the activation of some potential nucleation sites. The critical size of nucleus for recrystallization, r*, is given by

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 27A, OCTOBER 1996~2947

1-3000

Density of nucl. sites [ N1 t N2 N2 �9 N1 I H1 H2

Extent of cold work I [ / i IOOG H2 > H1 I

" 100 gg

e .

| ! i

(f/r) t (t/r) (f/r)crit

Fig. 1--Effect of second-phase particles' volume fraction to size ratio (fir) on gram size in a cold-worked and annealed metal (schematic). Region A final grain size controlled by Dgg, B- -by Dr,c, and C--stable subgrain structure.tin

Table I. Compositions, in Weight Percent, of Alloys (with AI Constituting the Remainder)

Alloy Mg Mn Cr Zr Fe Si

1 (5083) 4.93 0.81 0.19 0.001 0.084 0.034 4 (5083 + Zr) 4.66 0.83 0.16 0 .19 0.082 0.035 6 (A1-Mg) 4.74 0.001 0.00 0 .00 0.088 0.032 7 (5083 + Mn) 4.80 1.60 0.18 0.002 0.077 0.033 9 (5083 + Mn/Zr) 4.70 1.57 0.18 0.2 0.086 0.05

r* = Ky/(Pa - P,) [21

where y is the surface energy, Pd is the driving pressure for recrystallization, and K is a geometric constant. With increasing Pz, fewer nucleation sites get activated and Drec increases, as given by

Dre c = Ntot - '/3 [3]

where N,o t is the total density of activated nucleation sites. Once recrystallization is complete, particle drag on grain boundaries tends to restrict their migration and can thus enable a fine grain size to be retained. If the matrix contains a sufficiently large volume fraction of fine particles, such that P~ exceeds Pa, static recrystallization can be sup- pressed, leading to a stable subgrain structure. This has been the basis for a few semiempirical models proposed recently,t~4aSl as summarized in Figure 1 (Reference 15).

Figure 1 illustrates the competitive nature of the effects of the particle concentration and size, the (fir) ratio. In region A, the number of grains per unit volume, determined by the Zener drag, is smaller than that determined by the nucleation density and, thus, grain coarsening occurs after recrystallization. Grain growth occurs to a size, Dgg, given by

Dgg = a(r/f) [4]

where a is a constant of the order of unity. In region B, Dro c is larger than the size that can be stabilized by the particle distribution, (i.e., Dgg) and so there is no grain

growth after recrystallization. The value (fir), marks the transition between the two stages of the final grain size being dependent on grain growth and on the as-recrystallized size. Above a critical value, (fir)rot, Pz exceeds Pal, leading to subgrain stabilization (region C). The stabilized subgrain structure has been found to be a prerequisite for the continuous dynamic recrystallization during SPD, a mechanism that has been shown to lead to relatively high SPD rates.*

*There is an assumption here that particle dispersion remains constant throughout the testing period. Ostwald ripening of particles could lead to a decrease in (fir) and increase in Dee. Also, low mobility of grain boundaries might lead to smaller measured grain sizes than calculated, especially at shorter time intervals.

In this study, additions of particle-forming elements like Mn and Zr were made to the base alloy, along with TMP modifications, in order to vary the (fir) ratio and to assess the effect on recrystallization and grain size. Of interest were the possibilities for subgrain stabilization, or further grain refinement in a fully recrystallized condition. Recrys- tallization studies were conducted at temperatures of 0.5 Tm and above, because this is required for SPD.

II. EXPERIMENTAL PROCEDURE

A. Materials

Alloys of compositions, as shown in Table I, were sup- plied in the form of 12.7-ram plates, and cast and rolled by the Kaiser Aluminum Center for Technology (Pleasanton, CA). The alloys were given a precipitation anneal of 24 hours at 400 ~ after direct chill casting, and then hot rolled at the same constant temperature. This represents a modification of the usual TMP for this alloy, viz. a precipitation anneal of 10 hours at 510 ~ followed by hot rolling at the same temperature. This was done with a view to reduce the sizes of the intermetallic phase particles that form during the precipitation anneal.

Alloy 1 represents the base 5083 alloy (A1-4.75Mg- 0.8Mn), alloy 4 has extra Zr, alloy 7 has extra Mn, alloy 9 has extra Mn and Zr, while alloy 6 essentially has no alloying additions other than Mg. The Mg present in all alloys is added for solid solution strengthening, while Mn, Cr, and Zr are added to form intermetallic precipitates such as A16Mn and A13Zr. The elements Fe and Si, present as impurities, also form similar intermetallic particles. The Fe and Si contents here are somewhat lower than in commercial grade alloys, where they usually are about 0.4 pct and 0.3 pct, respectively. This was done to reduce the amount of large intermetallic particles that these elements form with A1 upon casting, which can contribute to cavitation and premature failure during SPD. All alloys also had less than 0.002 pct Cu, 0.02 pct Zn, 0.03 pct Ti, and 0.02 pct Ni.

B. Thermomechanical Processing

The alloys were rolled from 12.7 mm to 1.8-mm-thick sheets (true thickness strain of 1.97), by warm rolling at 150 ~ after a prior aging treatment at 350 ~ for 48 hours. The intent of the aging treatment was to produce a fine dispersion of particles that could pin the dislocations gen- erated during rolling, and thus influence recrystallization in

2948--VOLUME 27A, OCTOBER 1996 METALLURGICAL AND MATERIALS TRANSACTIONS A

the sheet alloys. Aging had to be done prior to rolling, because it was observed that rolling followed by aging re- suited in rapid recrystallization during the aging process itself. The aging temperature of 350 ~ was chosen to be above the Mg solvus temperature (300 ~ and prevent the precipitation of large AI3Mg 2 particles, which could hinder the suppression of static recrystallization, by serving as nucleation sites. Higher aging temperatures would result in fewer and coarser particles. Warm rolling was done to prevent cracking in the highly alloyed materials.

C. Recrystallization Studies

The recrystallization characteristics of all alloys were studied by annealing them over a range of temperatures from 300 ~ to 525 ~ with annealing times of 10 seconds to 24 hours. The shorter annealing operations (<1200 seconds) were carried out in a salt bath furnace. Temperature fluctuations (up to 15 ~ did not permit longer anneals to be similarly carded out, and an air furnace was used. The extent of recrystallization was determined by a combination of hardness measurements and metallographic examination. Hardness testing was conducted using a Rockwell testing machine and Rockwell B values were determined by av- eraging at least five test values. (It may be noted that the air furnace results in a slower heat-up rate of 75 ~ to 100 ~ per minute, compared to 100 ~ to 150 ~ per second for the salt bath furnace, which could affect recrystallization. However, as will be seen later, the differences in recrystallization between alloys 1, 6, and 7 vs 4 and 9 are quite obvious by 1200 seconds, and the overall results and discussion are not significantly altered.)

D. Metal lography

The primary objective of metallography was to confirm recrystallization and measure the as-recrystallized grain size, Drec, and establish if grain growth had occurred. Me- tallographic samples were given a pre-etch anneal at 180 ~ for 20 hours and etched with 10 pct H3PO4 at 50 ~ The objective of the preceding anneal was to precipitate the A13Mg2 particles along grain boundaries, which could be delineated by the subsequent etch. The microstructure was viewed in the plane containing the longitudinal and short transverse directions, and grain sizes were determined by the ASTM E112 standard intercept method using three con- centric circles.

E. Particle Content (17r)

Particle sizes and distribution were determined by scan- ning electron microscopy (SEM) and transmission electron microscopy (TEM) studies and also estimated from pub- lished literature, v5-~81 Identification of the particles was supported using the technique of X-ray fluorescence mapping in the SEM. A JEOL* JSM6400 SEM, operating at 10 kV,

*JEOL is a trademark of Japan Electron Optics Ltd., Tokyo.

was used for analyzing the K~ or L~ spectral energies of each element. X-ray fluorescence analysis gave not only the sizes of particles, but also information on the qualitative compositions of the particles. An estimate of sizes and com-

position of the bigger particles (> 1 ~m) was made using a computer image analysis method. The TEM studies were conducted (using a Hitachi H600 TEM, operating at 100 kV, or a JEOL 1200 operating at 120 kV) to obtain the size and distribution of smaller particles.

The volume fraction estimates were established for 350 ~ which was the aging temperature used. Weight fractions of intermetallic phases were obtained from the correspond- ing binary phase diagrams, such as A1-Mn and A1-Zr. tI3,~41 Approximate values of densities of different phases were used to arrive at the volume fractions. The f values, thus obtained, were really more of an upper bound or a maxi- mum achievable value o f f Such a single value of f was more convenient in assigning a numerical value to (fir), especially in the light of an inhomogenity in the distribution of particles across the matrix. This inhomogenity was also the reason why the preceding method was used for estimating f as opposed to, say, local estimates of the areal fraction from TEM micrographs. There is also some ques- tion about the applicability of data from binary phase diagrams to these ternary or quartemary (or higher) engineering alloys. It has been noted that the addition of Mg lowers the solid solubility of Mn in A1, t191 which could lead to a higher volume fraction of A16Mn particles than would be estimated from binary diagrams. However, the data available at the temperature and composition of interest is rather limited, and the uncertainty in values might well be in the range of the error listed in Table IV. So, as in a previous study, [14] this simplification was used.

III. RESULTS

A. Recrystallization Kinetics

On the basis of hardness values and supported by metallographic examination, a recrystallization parameter, /, was calculated, which is assumed to relate linearly to the fraction of material that had recrystallized, an approach pre- viously utilized, t2q

I = (H o - I-I,)/(Ho - /-/10o) [5]

where

H, /40

= hardness of the sample, after annealing; = hardness value of the unrecrystallized material;

and //loo = hardness of the fully recrystallized material.

Optical microscopy was used to monitor the progress of recrystallization, determine the 100 pct recrystallized condition, and provide qualitative support for the extent of recrystallization. Figure 2 shows typical micrographs of partially and completely recrystallized structures. Figure 3 shows the variation in hardness in alloys of different compositions following a recrystallization anneal at 300 ~ Al- loy 6 with no additions of Mn or Zr consistently has a lower hardness value and also exhibits the most rapid decrease in hardness. Alloy 9, on the other extreme, with Zr and extra Mn, exhibits the highest hardness, which remains quite stable over two time decades from 103 to 105 seconds.

Graphs of the I parameter, as a function of time for these alloys, gave the recrystallization kinetics and enabled the elucidation of the effects of composition (Figure 4) and recrystallization temperature (Figure 5).

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 27A, OCTOBER 1996--2949

Fig. 2--Partially and fully recrystallized microstructures: (a) partial recrystallization in alloy 4, following 1 h anneal at 300 ~ and (b) complete recrystallization in alloy 7, following 10 mm anneal at 525 ~

90

8O ~...__~~ .~

70 "--',~ ~ m ~ ,

\\\ o 60 ~ - " ' ~

,D ee 50

W w Q c s _

M Z

40

30 l ....

20

10 i ~

0 10 0 10 1 10 2 10 3 10 4 10

Annealing Time (sec) Fig. 3--Hardness variation with time, in different alloys, upon anneahng at 300 ~

.......... ~ ........ 6 (AI-Mg)

~- 1(5083)

= 4 (5083+Zr)

o 7 (5083+Mn)

......... r ...... 9 (5083+Mn+Zr)

The recrystallization kinetics at 300 ~ of the various alloys are shown in Figure 4, which is directly derived from Figure 3. Alloy 6 recrystallizes completely within 200 seconds at 300 ~ Alloys 1 and 7, with Mn additions, also exhibit complete recrystallization, although with slower kinetics. However, those alloys containing Zr (4 and 9) exhibit a tendency to arrest the full recrystallization at 50 to 60 pct, and this arrested recrystallization appears to be sta-

ble for at least 105 s. There is no clear evidence that the addition of extra Mn, over and above the commercial base alloy (in the absence of Zr), affected the recrystallization kinetics within the range evaluated, because alloy 7 containing 1.6Mn exhibited full recrystallization with about the same kinetics as alloy 1 contains 0.8 pct. Alloy 9 was the most resistant to recrystallization at 300 ~ but seems to recrystallize with characteristics not too different from alloy

2950~VOLUME 27A, OCTOBER 1996 METALLURGICAL AND MATERIALS TRANSACTIONS A

1.0'

0.8 '

Q N

m m 0.6 '

O Q L -

r 0 .4 ' O

o 10

0 .2 '

0.0'

/ f

~ - - I ~ 1 I I I I I l l �9 l I l s I l l

0 0 01 10 2

f

I !

l

r

7, !

' e - /

!

: A : t l

I

I l 1 I

a ~

/ , a

/ - , , , , , ,

I

r

!

!/ i I ~ / s ! �9 I I / i / "

/ I j l l �9

"1 i l ~ / l I

/ i l l ~

/ � 9 t / d_?t/ J A . . . . . . . . .

i / rY"

B - &

A I

- J-A-a

q

�9 �9 �9 i � 9 1 4 9 �9 �9 | | 1 � 9 1 4 9

.10 3 10 4 10 5

�9 6 (AI-Mg)

13 1 (5083)

o 7 (5083+Mn)

& 4 (5083+Zr)

�9 9 (5083+Mn+Zr)

Annealing Time (sec)

Fig. 4---Recrystallization kinetics, measured at 300 ~ of alloys of different compositions (derived from Fig. 3).

A 'u

O

e- O m

o

1.0

08

06

/ / ~ 1 1

i P/"

/

t / C - " - U I,; 1 / u

i I

I

0.4 / / /

t / /

; , i 0.2 J iu /

a /

D

O0 - - ~ ' 10 0 101 10 2 10 3 10 4 10 5

Annealing time (aec)

Fig. 5--Recrystalhzatlon kinetics of alloy 9 (5083 + Mn + Zr) at two different temperatures. Rapid recrystallization occurs at 525 ~

4, which has a similar Zr content but lower Mn content. The partial recrystallization in alloys 4 and 9 suggests a nonuniformity in the particle distribution, with suppression of static recrystallization occurring in regions where (fir) exceeds QTr)cr, t. Therefore, regions with a high local concentration of fine particles, or alternately with very few large particles, could remain unrecrystallized. Figure 2(a), which shows regions of fully recrystallized microstructures with adjoining bands of unrecrystallized material in alloy 4, appears to support the previous statement.

It must be noted that local variations in Pd could cause variations in (fir),, itself; i.e., in the unrecrystallized regions, Pd itself could be lower in addition to P: being higher. Upon rolling, Pd might vary from grain to grain,

depending on the initial grain orientation. This possibility was examined as follows. The average grain size in the as- received material (12.7-mm thick), measured for alloy 4 after a 525 ~ h anneal, was 23.4/xm, giving an average area of 430/xm 2 per grain. While some of the unrecrystallized regions in Figure 2(a) have an area approximately equal to the above area, larger areas in excess of 800/zm 2 have also been observed, indicating the presence of at least two original grains in that region. The former could suggest that Pa varies from grain to grain, but the latter suggests that some other factors also affect recrystallization (viz. inhomogeneity in (fir) and resultant P.). So, while the possibility of an inhomogeneity in the stored energy cannot be altogether ruled out (at least in the smaller unrecrystallized regions), the inhomogeneity in particle densities appears to contribute more strongly.

Alloy 9 does exhibit rapid and complete recrystallization without an arrest plateau at 525 ~ as shown in Figure 5, where the 300 ~ recrystallization properties are also shown for comparison. At this higher temperature, complete recrystallization occurs within about 100 to 500 seconds. In a typical SPF operation at 525 ~ with this alloy, the heat- up and temperature stabilization times are usually at least 30 minutes. So, it can be concluded that even alloy 9, with its extra Mn and Zr additions, would be in a statically recrystallized condition at the start of SPF.

B. Recrystallized Microstructure

Grain size measurements after different holding times at 525 ~ are shown in Table II. The term Droc is the grain size measured in the as-recrystallized condition. It was measured after a 10-second anneal for alloy 6, while other


Table II. Grain Sizes in Alloys, after Recrystallizing Anneal at 525 ~ for 10 Minutes (Dr,0) (10 Seconds for Alloy 6) and Prolonged Exposure (D after 1 and 24 Hours)

Alloy

Aspect Ratio in DT~ c (p~m) As-Recrystallized D (/~m) D (/zm)

(525 ~ min) Grains (525 ~ h) (525 ~ h)

1 5.80 2.25 6.62 7.56 4 5.94 2.86 6.03 6.37 6 8.69 (10 s) 1 40.86 61.69 7 6.11 1.65 6.17 7.02 9 6.05 2.35 5.94 6.54

Fig. ~ a ) and (b) As-recrystallized grain sizes D~c in alloys 6 (A1-4.75Mg) and 9 (AI-4.75Mg-l.6Mn-0.2Zr), respectively, upon recrystalhzation anneal at 525 ~ D~c measured after 10 s anneal for 6 and 10 min anneal for 9. (c) and (d) Grain sizes in alloys 6 and 9, after a prolonged exposure of 1 h. Alloy 6 shows pronounced grain coarsening upon longer exposures, while 9 is quite stable.

alloys were annealed for 10 minutes. This is because alloy 6 was found to recrystallize very rapidly and expected to

grow rapidly as well. Grain growth was measured after 1 and 24 hours. Figure 6 shows typical optical micrographs


of the initial as-recrystallized grains in alloys 6 and 9 and of their growth following 1-hour exposures.

As seen from the aspect ratios presented in Table II, alloys 1, 4, 7, and 9, with higher amounts of alloying additions of Mn and Zr, have elongated grains, while alloy 6 has an equiaxed grain structure, which is retained as growth occurs. Grain size results from Table II are graphically rep- resented in Figure 7. The addition of alloying elements seems to stabilize the grain structure, as seen by the large amount of grain growth in alloy 6, as opposed to the other alloys.

C. Particle Sizes and Distribution

Electron microscopy studies revealed particles containing Mn and Cr, which were probably A16(Mn,Cr ) and those containing Zr, which appear to be A13Zr. Some of the larger Mn-bearing particles (>2/zm) also contained Fe. Figure 8 shows SEM micrographs of the A16(Mn,Cr ) precipitate structure in the 5083 base alloy (alloy 1) and 5083 with additional Mn (alloy 7). The higher Mn addition in alloy 7 resulted in only a small increase in the density of submicrometer particles, but the density of large precipitates (>2 /zm) increased almost four times, from 0.0017/xm 2 in alloy 1 to 0.0064/~m -z in alloy 7.

The dispersion of finer particles is shown in the TEM micrographs of alloy 1 (5083) and alloy 4 (5083 with Zr) in Figure 9. A13Zr particles present in alloy 4 are much finer than A16(Mn,Cr ) particles, which are present in both alloys. While the A16(Mn,Cr ) particles were incoherent, both coherent and incoherent A13Zr particles were observed. A13Zr particles were in the form of submicrometer particles, whereas the Mn-containing particles ranged from 0.04 to 7.1 /zm. Table III shows a compilation of the sizes of the particles, based on the preceding electron microscopy studies and literature, t~6,~7,tsl Volume fractions of the particles in different alloys, estimated from phase diagrams, are shown in Table IV. These data are utilized in Section IV to estimate a total (fir) for the different alloys.

IV. DISCUSSION

A. Recrystallization Kinetics

The preceding results indicate possible explanations for the observations made from Figures 4 through 6. As seen from Figure 8, the addition of excess Mn to 5083 (alloy 7) leads to an increase in the number of nucleation sites, due to the fourfold increase in the density of particles >2/xm, while there is only a small increase in the submicrometer particle density. The observation that an increase in Mn content alone (from 0.8 pct in alloy 1 to 1.6 pct in 7) does not inhibit recrystallization very much is consistent with this result (Figure 4). Alloy 9 exhibits a higher resistance to recrystallization than 7, because Zr additions are very effective in inhibiting recrystallization because of their smaller size. Thus, despite being present in lower propor- tion than Mn, they impose a higher drag.

The preceding results, such as the rapid recrystallization in alloys 1 and 7, as opposed to 4 and 9, reinforce general ideas of the role of the fine Zr-containing particles in suppression of static recrystallization by restricting the mobility of moving boundaries. The effect of Mn additions is a little

100

o

g

D rec ~

1 . . . . . . . , i . . . . . . . . i . . . . . . . . | . . . . . . . . | . . . . . .

10 0 101 102 10 3 104

Annealing Time (sac) 10 5

Fig. 7- -Gram sizes in different alloys upon anneahng at 525 ~ Alloy 6, with no Mn or Zr additions, exhxbits significant grain growth with time.

different. While an increase in Mn content from 0 (alloy 6) to 0.8 pct (alloy 1) helps reduce recrystallization kinetics, it clearly does not lead to its complete suppression. A further increase in the Mn content to 1.6 pct (alloy 7), while producing a slight increase in the number of finer particles, results in a much larger increase in coarser particles. The latter effect appears to negate any beneficial effect from the former in terms of suppressing static recrystallization.

B. Particle (f/r) Values and Relation to Grain Size

Results of the microstructural analysis clearly indicate a bimodal distribution of particles (Table III). A simplifying assumption, that the predominant drag pressure is due to fine particles and that of the larger particles is almost neg- ligible, can help make approximate (fir) calculations for the purpose of quantifying Figure 1. This assumption appears to be quite reasonable, as reported by Welt and Austin. [~41

The effect of the large particles, however, cannot always be overlooked, as they do have an indirect influence on recrystallization by varying the number of potential nucleation sites N,. (How significantly N, is varied depends on the relative density of other potential nucleation sites like triple points, shear bands, etc.) Using the volume fractions data in Table IV, taking conservative estimates of sizes rA,~(Mn,c,~ = 0.2 /xm and rA~3z, = 0.025 /zm (Table III), and following the Wert and AustinI~"] method of cumulative estimation of (fr), we can get approximate values of (fir) for different alloys, as in Table V. The preceding method, in this case, becomes

(f/r)tota, = (f/F)AI6(Mn,Cr) d- 2(f/r)A,,Zr [6]

The multiplicative factor 2 for the A13Zr particles arises from earlier work by Nes et al. t~31 showing that coherent particles exert twice the drag of incoherent particles. In our study, although some large A13Zr particles (with sizes >0.05/zm) were found to be incoherent, for simplicity, we treat all Zr-bearing particles as coherent.

The estimated (fir) values, together with the measured grain sizes from Table V, can be utilized in the graph of D vs (fir) (Figure 10). Grain size measurements indicate a very pronounced growth of grains with time, in the case of


Fig. 8---(a) and (b) SEM mlcrographs showing the AI6(Mn,Cr ) precipitate structure in the 5083 base alloy (alloy 1) and 5083 with additional Mn (alloy 7), respectively. Addition of Mn results in a fourfold increase in the number of large (>2/zm) precipitates.

Fig. 9--(a) and (b) TEM micrographs showing the dispersion of second-phase particles in the 5083 base alloy (alloy 1) and 5083 + Zr alloy (alloy 4), respectively. Larger particles of A16(Mn,Cr), along with finer Al3Zr particles.

alloy 6, with a smaller growth was observed in alloy 1, and the absence of significant growth, in the case of alloys 7, 4, and 9. Also shown is a theoretically predicted limiting grain growth size, from Eq. [4], using a of unity. All alloys

possess an initial fine-grained structure, which can be at- tributed to the high degree of cold work ( - 8 6 pct). But clearly, in alloy 6 with no Mn and Zr additions (the only particles possibly present being those containing the trace

2954~VOLUME 27A, OCTOBER 1996 METALLURGICAL AND MATERIALS TRANSACTIONS A

Table III. Size Distribution of Particles, Obtained from SEM and TEM Studies and Estimates from References 6*

and 16 through 18"*

Diameter Diameter Coarse Fine

Particles Particles Particle (/Lm) (/zm)

A16(Mn,Cr) 1.5 to 7.1 0.04 to 0.2 A13Zr 0.05 to 0.1" 0.01 to 0.025**

Table IV. Equilibrium Volume Fractions ( f ) of Various Particles in Modified 5083 Al Alloy at 350 ~ (Estimated

from Phase Diagrams, References 19 and 20)*

A16(Mn,Cr) A13Zr Alloy (Error +__ 0.0015) (Error _+ 0.0005)

1 0.026 - - 4 0.026 0.0025 6 - - - - 7 0.052 - - 9 0.052 0.0025

*Source of estimated error: least count of measuring ruler.

Table V. Estimation of Total (f/r) for Alloys at 350 ~

Alloy Total (fir) 1 0.13 4 0.33 6 0.00 7 0.26 9 0.46

elements), the limited particle dispersion cannot stabilize this initial fine-grained structure. This is seen in the grains in this alloy, growing from Dre ~ (which is about 9 /zm) to Dgg, further growth being restricted by particle dispersion. The continued grain growth in this alloy between the 1- and 24-hour annealing periods could indicate that the equilibrium Dgg value in this alloy has not been attained after 1 hour, due to relatively slow grain mobility. Alternately, there could be Ostwald ripening or dissolution of the precipitates. In contrast to alloy 6, the grain size in alloy 9 is stable at the level indicated by Dr,r Grain size trends, noted by Vetrano et al.,t6I seem to support this view.

It appears, therefore, that with an increase in (fir), there is a transition from grain growth controlled grain sizes (i.e., Dgg controlled), as in alloy 6, to one that is controlled by the as-recrystallized grain size (D J , as in alloy 9. The grain growth observed in alloys 1 and 7 is quite small (being less than an order of magnitude), which could imply that the particle contents in these alloys are quite close to the (fir), values for the alloys. A rough sketch of the Dr~o profile is shown, based on these considerations. Additional data points at higher (fir) values could help obtain better experimental quantification, but this rough curve is useful for illustrative purposes.

A point of interest is why there is not much change in the measured as-recrystallized grain size, D~r with (fir). The as-recrystallized sizes in alloys 1, 4, 7, and 9 are almost the same, about 6/xm, indicating a nearly constant density of nucleation sites in all of the heavily cold-rolled materials,

when recrystallized under these conditions. The ratio of the density of as-recrystallized grains to the density of particles larger than 2 /zm for the alloy 1 was found to be 22.2, while that for alloy 7 with additional Mn was 5.3. This indicates that there are many more nucleation sites in operation, in addition to the large particles in alloy 1, and, to a lesser extent, in alloy 7. So, while large particles can influence the nucleation of recrystallization, they are not the only sites where nucleation can occur. This is especially true of alloys that have been heavily cold rolled and/or recrystallized at higher temperatures. It is, of course, possible that a large particle might nucleate more than one grain, but in view of such high ratios, it is quite reasonable to expect other nucleation sites being operative.

C. Estimation o f (f/r) .... and Suggested Alloy Modifications

From the preceding results, it is apparent that the addition of Zr and Mn has resulted in an increase in (fir), but this has not been sufficient to prevent the onset of static recrystallization. In this section, an estimation of the critical value of (fir) for subgrain stabilization (at a size of 1 /zm) has been attempted.

(a) One approach toward estimating (fir)on,, similar to that adopted by Nes,tl8j is to consider the equation

Dgg = a(r/f) [4]

where a is a geometric constant of the order of unity (=4/3 in Zener's original analysis) and Dgg is now the size of the subgrain that would be stabilized by the particle dispersion. Therefore, to stabilize a subgrain of 1 /zm, assuming a = 1,

( f /r)cn,- 1 /xm -l [7]

(b) A more rigorous method is to determine the value of (fir) at which the driving pressure for recrystallization, Pa, equals the Zener drag pressure, ~ , which is the desired (fir)r Wert and Austin t141 have reported measurements of P~ as a function of rolling strain for the A1-0.38 pct Zr system. Using these values as a first approximation, for the rolling strain of 1.97 employed in this study,

Pd = 0.46 MPa = P_ = 1.5y(f/r)~,t [8]

For 3' = 0.3 J/m2, ~131

(f/r)~,, = 1.022 /zm -1 [9]

Assuming a value of (fir)r , of 1 /zm l, the alloy compositions that could lead to this value can be estimated. The following assumptions have been made.

(a) The alloy has been subjected to an identical TMP, in- cluding the 350 ~ h preroll aging treatment.

(b) The sizes of the Mn and Zr intermetallic particles are assumed to be the same as those used in (fir) estimates reported in Table VI, viz. rA~,tM,,Cr.Fe ~ = 0.2 /xm and FAI3Z r = 0.025 tzm.

Based on volume fraction data calculated at 350 ~ and using the preceding particle sizes, the alloys given in Table VI have been calculated to have an (flr),o , equal to the calculated (fir) .... of 1 /zm i.

METALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 27A, OCTOBER 1996--2955

100

1 / [ ~'~= i D rec (sketch) /

r \ ~ / / 0 L . ,

.0 o D rec

E 6 ~kk / , , / n D(lh)

"wN =1 10 }" .~.~ g ~ . ~ ,,," " "~J " .......... '= DD(24h)gg

- - . . . . Drec(sketch) " - - 9 Ig

1 , I , I �9 i . , . - - - . - . . ~

0.0 0.2 0.4 0.6 0.8 1.0 (f/r)

Fig. 10--Furu graph, showing grain sizes of alloys 1, 4, 6, 7, and 9, as a function of their (fir), after 10 min, 1 h, and 24 h recrystallization anneals at 525 ~ (from data in Tables II and V). (Droo for alloy 6 was done after a shorter anneal of 10 s.) Also overlaid is the theoretically predicted limiting grain growth size, based on Eq. [4], and a rough sketch of the Drr profile.

It seems quite unlikely that such fine and uniform particle dispersions can be achieved in alloys of some of the above compositions, especially those involving the higher Mn contents. The results of this study bear out the relative coarsening of the Mn-containing particle structure, upon increasing Mn content (alloy 7 as opposed to alloy I). So, while the success of achieving the desired fine particle dispersion throughout the matrix in these alloys, by the normal ingot casting process, is slight, there is some evidence to indicate the possibility of subgrain stabilization by other processing methods. Higashi et al. 171 have reported a subgrain structure of about 1 /zm in a A1-5Mg-0.8Zr alloy, formed by PM, in which dynamic recrystallization was observed upon SPD at 500 ~ This alloy composition appears quite consistent with alloy I in Table VI.

D. Limitations o f Present Approach

The model developed by Furu et al.[15] can be used to predict conditions, resulting in complete suppression of static recrystallization. The authors assume a uniform distribution of particles in the matrix. This is not always valid in real situations, where a nonhomogeneous distribution of particles may be present. In that case, some regions may have a high enough local value of (fir) to prevent static recrystallization, while others may not, leading to partial recrystallization, as seen in alloys 4 and 9 at 300 ~ (Figure 4).

The "Furu graph" (Figures 1 and 10) is helpful in un- derstanding results of the study and suggests that the (fir) ratio, while increased above that of the 5083 alloy on add- ing Mn and Zr, was not large enough to fully suppress the

Table VI. Alloy Concentrations in the Base Alloy Al- 4.75Mg Estimated for a 1-pm Subgrain Stabilization

Alloy Wt Pct Mn Wt Pct Zr

I 0.00 0.66 II 0.80 0.58 III 1.60 0.49 IV 4.20 0.20 v 5.20 0.00

onset of recrystallization for the heavily worked materials. While the magnitude of (fir) is helpful, the determination of a single value of (fir), based on a single value o f f and an average value of r, as has been done here, has its limitations. The aspect of the microstructure not reflected by a total (fir) is the particle-size distribution. It is expected that larger particles (greater than 1 to 2/xm), which would be large with respect to the subgrain size, would tend to stim- ulate recrystallization. The finer particles (less than 1 to 0.5 /~m) may act to inhibit recrystallization by pinning the subgrain structure. When both large and small particles are present, the situation is more complex, as there would be competition between these two effects. Regions around some large particles may have insufficient fine particles to prevent nucleation of recrystallization. Conversely, there may be regions devoid of large particles, or in which sufficient fine particle concentrations are present, so as to inhibit nucleation of recrystallization. In the case where both of these effects may be present, one may observe rapid onset of recrystallization but inhibition of full recrystalli-


zation. The results of the research reported here for alloys 4 and 9 are consistent with this concept.

V. CONCLUSIONS

An increase in the Mn content alone from 0.8 to 1.6 pct did not impede recrystallization due to the formation of big precipitates. The Mn increase, when coupled with the addition of 0.2 pct Zr, helped suppress static recrystallization, at least partially. At higher temperatures, where complete recrystallization occurred, the increase helped stabilize the grain structure. An estimate of the critical (fir) ratio required to stabilize a 1-/zm subgrain structure was found to correspond to a significant increase in the desired second- phase particle density and a suitable modification in the chemical composition. In practice, it seems unlikely that such fine particle distributions could be achieved in this alloy system by ingot metallurgy. The model of Furu et al.t~sJ was adopted to understand the grain size results obtained on recrystallization, in terms of (fir) of particles. While their model is a useful qualitative tool, it was found that the particle size distribution and inhomogeneity in (fir) across the metal sample tends to restrict its applicability.

ACKNOWLEDGMENTS

This research was supported by a grant from Pacific Northwest National Laboratory. The work of JSV was supported by the Materials Science Division, Office of Basic Energy Sciences, United States Department of Energy (DOE), under Contract No. DE-AC06-76RLO 1830. The authors would also like to thank Curt Lavender and Mark Smith, Pacific Northwest National Laboratory, for helpful discussions.

REFERENCES

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Bruemmer: Scripta Metall Mater., 1994, vol. 30, pp. 565-86. 4. H. Iwasaki, K. Higashi, S. Tanimura, T. Komatubara, and S. Hayami:

Superplasttctty in Advanced Materials, ICSAM '91, S. Hori, M. Tokizane, and N. Furushiro, eds., Japan Society for Research on Superplasticity (JSRS), Osaka, Japan, 1991, pp. 447-52.

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16. T.R. McNelley and S.J. Hales: Superplasticity in Aerospace 11, T.R. McNelley and C.H. Heikkenen, eds., TMS, Warrendale, PA, 1990, pp. 207-22.

17. T.R. McNelley and P.N. Kalu: Superplasticity in Advanced Materials, ICSAM '91, S. Hori, M. Tokizane, and N. Furushiro, eds., Japan Society for Research on Superplasticity (JSRS), Osaka, Japan, 1991, pp. 413-21.

18. E. Nes: Met. Sci., 1979, Mar.-Apr., pp. 211-15. 19. Metals Handbook, 8th ed., ASM, Metals Park, OH, 1973, Butterworth

and Co., London, vol. 8, pp. 258-65 and 342-45. 20. L.F. Mondolfo: Aluminum Alloys--Structure and Properties,

Butterworth and Co., London, 1976, pp. 250, 324, 338, and 413. 21. E.S. Puchi and G. Rojas: Mater. Sci. Forum, 1993, vols. 113-115,

pp. 509-14.

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