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

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  • 1. Effects of Alloy Modification and ThermomechanicalProcessing 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 re- crystallization 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 par- ticle 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. tructure, the higher the strain rate for superplastic I. INTRODUCTION deformation (SPD). Two types of microstructures are be- lieved to be suitable for improved SD rates: (1) stabilized THE A1-Mg-Mn alloy 5083 is of interest for superplastic subgrain structure leading to continuous dynamic recrystal-forming (SPF) because of its moderate strength, corrosion lization during SPD and (2) statically recrystallizedresistance, weldability, and because it has been shown to equiaxed microstructure with grain sizes less than about 5be sufficiently superplastic to permit SPF of components, t~- /xm, which is stable at temperatures greater than 0.5Tin (in61 Elongations of over 300 pctt2.3] have been achieved using degrees kelvin). The former has been demonstrated in thecommercial grades of the 5083 alloy. Development of SPF- 8xxx A1-Li-Zd 891 and the Supral 2xxx t~~ series of alloysgrade 5083 alloys, chiefly by reducing the Fe and Si im- (by ingot casting), but there have not been many such stud-purity content, has led to higher elongations of up to 630 ies with the 5083 series. Almost all studies reported for thispct. H,51Other modifications have involved additions of Cu, alloy have been with (2), i.e., statically recrystallized mi-Zr, and Mn.t4.61 Vetrano et al.t3] also studied the effect of crostructures, with grain sizes of about 8 to 10/zm.thermomechanical processing (TMP), aimed at reducing Recrystallization characteristics and resulting grain sizesgrain sizes, on superplastic elongation in commercial al- in A1 alloys, arising out of a particular TMP, are related toloys. This study, which focused on the effect of cold rolling the amount and distribution of second-phase particles in thereductions, concluded that thickness reductions of at least matrix. Hard, nondeformable particles, larger than a critical5:1 gave the best superplastic tensile elongations. size, can serve as sites for recrystallization and, thus, influ- In most cases, the strain rates necessary to obtain ac- ence the as-recrystallized grain size Droc. For most AI al-ceptable ductility are in the range of 2 • 10 4 S-I to 4 • loys, this critical particle size has been observed to be10 -4 S- I , requiring relatively long forming times, of the around 1 to 2 /xm.[~~,~21An increase in the density of suchorder of 30 minutes to a few hours. There is, therefore, nucleation sites leads to finer recrystallized grains. How-interest in increasing the superplastic strain rate of this type ever, another aspect of the effect of particles, viz. the Zenerof alloy through further alloy development and TMP, while drag pressure, can sometimes lead to a contrary effect onretaining the low-cost casting methods for the material. recrystallization and growth, t13,14`~slParticles of all sizes ex-There has been some success in achieving elongations of ert a Zener drag pressure P~ on moving boundaries, restrict-500 to 600 pct at high strain rates of 10 -2 s -1 in the A1- ing their migration, where5Mg-0.8Zr alloy system processed by powder metallurgy(PM). t71 It is easier to obtain a finer dispersion of second- t 9_ = t9 y f / r [1]phase particles and finer grain sizes by PM than by ingotcasting, but the former is substantially more expensive. where y is the surface energy,/3 is a constant, a n d f a n d r It is well known that a primary factor in producing su- represent the volume fraction and size of the particle.perplasticity is a fine grain size, and the finer the micros- 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 K. KANNAN, Graduate Student, and C.H. HAMILTON, Professor, are leads to opposing effects on the final grain size. Drag pres-with the School of Mechanical and Materials Engineering, Washington sure on subgrain boundaries by the fine particles leads to aState University, Pullman, WA 99164-2920. J.S. VETRANO, SeniorResearch Scientist, is with the Structural Materials Research Section, coarser D .... by preventing the activation of some potentialPacific Northwest National Laboratory, Rlchland, WA 99352. nucleation sites. The critical size of nucleus for recrystal- Manuscript submitted October 30, 1995. lization, r*, is given byMETALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 27A, OCTOBER 1996~2947
  • 2. 1-3000 growth after recrystallization. The value (fir), marks the Density of nucl. sites [ N1t N2 transition between the two stages of the final grain size N2 9 N1 I H1 H2 being dependent on grain growth and on the as-recrystal- Extent of cold work [ I /i lized size. Above a critical value, (fir)rot, Pz exceeds Pal, IOOG H2 > H1 I 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 " 100 gg SPD rates.* e. *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, | ! i 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 (f/r) t (t/r) (f/r)crit modifications, in order to vary the (fir) ratio and to assess the effect on recrystallization and grain size. Of interestFig. 1--Effect of second-phase particles volume fraction to size ratio (fir) were the possibilities for subgrain stabilization, or furtheron gram size in a cold-worked and annealed metal (schematic). Region grain refinement in a fully recrystallized condition. Recrys-A final grain size controlled by Dgg, B--by Dr,c, and C--stable subgrain tallization studies were conducted at temperatures of 0.5 Tmstructure.tin and above, because this is required for SPD. Table I. Compositions, in Weight Percent, of Alloys (with AI Constituting the Remainder) II. EXPERIMENTAL PROCEDURE Alloy Mg Mn Cr Zr Fe Si A. Materials1 (5083) 4.93 0.81 0.19 0.001 0.084 0.0344 (5083 + Zr) 4.66 0.83 0.16 0.19 0.082 0.035 Alloys of compositions, as shown in Table I, were sup-6 (A1-Mg) 4.74 0.001 0.00 0.00 0.088 0.032 plied in the form of 12.7-ram plates, and cast and rolled by7 (5083 + Mn) 4.80 1.60 0.18 0.002 0.077 0.033 the Kaiser Aluminum Center for Technology (Pleasanton,9 (5083 + Mn/Zr) 4.70 1.57 0.18 0.2 0.086 0.05 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 modi- r* = Ky/(Pa - P,) [21 fication of the usual TMP for this alloy, viz. a precipitationwhere y is the surface energy, Pd is the driving pressure anneal of 10 hours at 510 ~ followed by hot rolling atfor recrystallization, and K is a geometric constant. With the same temperature. This was done with a view to reduceincreasing Pz, fewer nucleation sites get activated and Drec the sizes of the intermetallic phase particles that form dur-increases, as given by ing the precipitation anneal. Alloy 1 represents the base 5083 alloy (A1-4.75Mg- Dre c = Ntot -/3 [3] 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 al-where N,ot is the total density of activated nucleation sites. loying additions other than Mg. The Mg present in all al-Once recrystallization is complete, particle drag on grain loys is added for solid solution strengthening, while Mn,boundaries tends to restrict their migration and can thus Cr, and Zr are added to form intermetallic precipitates suchenable a fine grain size to be retained. If the matrix contains as A16Mn and A13Zr. The elements Fe and Si, present asa sufficiently large volume fraction of fine particles, such impurities, also form similar intermetallic particles. The Fethat P~ exceeds Pa, static recrystallization can be sup- and Si contents here are somewhat lower than in commer-pressed, leading to a stable subgrain structure. This has cial grade alloys, where they usually are about 0.4 pct andbeen the basis for a few semiempirical models proposed 0.3 pct, respectively. This was done to reduce the amountrecently,t~4aSl as summarized in Figure 1 (Reference 15). of large intermetallic particles that these elements form with Figure 1 illustrates the competitive nature of the effects A1 upon casting, which can contribute to cavitation andof the particle concentration and size, the (fir) ratio. In re- premature failure during SPD. All alloys also had less thangion A, the number of grains per unit volume, determined 0.002 pct Cu, 0.02 pct Zn, 0.03 pct Ti, and 0.02 pct Ni.by the Zener drag, is smaller than that determined by thenucleation density and, thus, grain coarsening occurs afterrecrystallization. Grain growth occurs to a size, Dgg, given B. Thermomechanical Processingby The alloys were rolled from 12.7 mm to 1.8-mm-thick Dgg = a(r/f) [4] sheets (true thickness strain of 1.97), by warm rolling at 150 ~ after a prior aging treatment at 350 ~ for 48 hours.where a is a constant of the order of unity. In region B, The intent of the aging treatment was to produce a fineDroc is larger than the size that can be stabilized by the dispersion of particles that could pin the dislocations gen-particle distribution, (i.e., Dgg) and so there is no grain erated during rolling, and thus influence recrystallization in2948--VOLUME 27A, OCTOBER 1996 METALLURGICAL AND MATERIALS TRANSACTIONS A
  • 3. the sheet alloys. Aging had to be done prior to rolling, position of the bigger particles (> 1 ~m) was made usingbecause it was observed that rolling followed by aging re- a computer image analysis method. The TEM studies weresuited in rapid recrystallization during the aging process conducted (using a Hitachi H600 TEM, operating at 100itself. The aging temperature of 350 ~ was chosen to be kV, or a JEOL 1200 operating at 120 kV) to obtain theabove the Mg solvus temperature (300 ~ and prevent the size and distribution of smaller particles.precipitation of large AI3Mg2 particles, which could hinder The volume fraction estimates were established for 350the suppression of static recrystallization, by serving as nu- ~ which was the aging temperature used. Weight fractionscleation sites. Higher aging temperatures would result in of intermetallic phases were obtained from the correspond-fewer and coarser particles. Warm rolling was done to pre- ing binary phase diagrams, such as A1-Mn and A1-Zr.tI3,~41vent cracking in the highly alloyed materials. 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-C. Recrystallization Studies mum achievable value o f f Such a single value of f was The recrystallization characteristics of all alloys were more convenient in assigning a numerical value to (fir),studied by annealing them over a range of temperatures especially in the light of an inhomogenity in the distributionfrom 300 ~ to 525 ~ with annealing times of 10 seconds of particles across the matrix. This inhomogenity was alsoto 24 hours. The shorter annealing operations (<1200 sec- the reason why the preceding method was used for esti-onds) were carried out in a salt bath furnace. Temperature mating f as opposed to, say, local estimates of the arealfluctuations (up to 15 ~ did not permit longer anneals to fraction from TEM micrographs. There is also some ques-be similarly carded out, and an air furnace was used. The tion about the applicability of data from binary phase dia-extent of recrystallization was determined by a combination grams to these ternary or quartemary (or higher)of hardness measurements and metallographic examination. engineering alloys. It has been noted that the addition ofHardness testing was conducted using a Rockwell testing Mg lowers the solid solubility of Mn in A1,t191which couldmachine and Rockwell B values were determined by av- lead to a higher volume fraction of A16Mn particles thaneraging at least five test values. (It may be noted that the would be estimated from binary diagrams. However, theair furnace results in a slower heat-up rate of 75 ~ to 100 data available at the temperature and composition of inter-~ per minute, compared to 100 ~ to 150 ~ per second est is rather limited, and the uncertainty in values mightfor the salt bath furnace, which could affect recrystalliza- well be in the range of the error listed in Table IV. So, astion. However, as will be seen later, the differences in re- in a previous study,[14] this simplification was used.crystallization between alloys 1, 6, and 7 vs 4 and 9 arequite obvious by 1200 seconds, and the overall results and III. RESULTSdiscussion are not significantly altered.) A. Recrystallization KineticsD. Metallography On the basis of hardness values and supported by me- tallographic examination, a recrystallization parameter, /, The primary objective of metallography was to confirm was calculated, which is assumed to relate linearly to therecrystallization and measure the as-recrystallized grain fraction of material that had recrystallized, an approach pre-size, Drec, and establish if grain growth had occurred. Me- viously utilized, t2qtallographic samples were given a pre-etch anneal at 180~ for 20 hours and etched with 10 pct H3PO4 at 50 ~ I = ( H o - I-I,)/(Ho - /-/10o) [5]The objective of the preceding anneal was to precipitate the whereA13Mg2 particles along grain boundaries, which could bedelineated by the subsequent etch. The microstructure was H, = hardness of the sample, after annealing;viewed in the plane containing the longitudinal and short /40 = hardness value of the unrecrystallized material;transverse directions, and grain sizes were determined by andthe ASTM E112 standard intercept method using three con- //loo = hardness of the fully recrystallized material.centric circles. Optical microscopy was used to monitor the progress of recrystallization, determine the 100 pct recrystallized con-E. Particle Content (17r) dition, and provide qualitative support for the extent of re- crystallization. Figure 2 shows typical micrographs of Particle sizes and distribution were determined by scan- partially and completely recrystallized structures. Figure 3ning electron microscopy (SEM) and transmission electron shows the variation in hardness in alloys of different com-microscopy (TEM) studies and also estimated from pub- positions following a recrystallization anneal at 300 ~ Al-lished literature,v5-~81Identification of the particles was sup- loy 6 with no additions of Mn or Zr consistently has a lowerported using the technique of X-ray fluorescence mapping hardness value and also exhibits the most rapid decrease inin the SEM. A JEOL* JSM6400 SEM, operating at 10 kV, hardness. Alloy 9, on the other extreme, with Zr and extra *JEOL is a trademark of Japan Electron Optics Ltd., Tokyo. Mn, exhibits the highest hardness, which remains quite sta- ble over two time decades from 103 to 105 seconds.was used for analyzing the K~ or L~ spectral energies of Graphs of the I parameter, as a function of time for theseeach element. X-ray fluorescence analysis gave not only the alloys, gave the recrystallization kinetics and enabled thesizes of particles, but also information on the qualitative elucidation of the effects of composition (Figure 4) andcompositions of the particles. An estimate of sizes and com- recrystallization temperature (Figure 5).METALLURGICAL AND MATERIALS TRANSACTIONSA VOLUME 27A, OCTOBER 1996--2949
  • 4. Fig. 2--Partially and fully recrystallized microstructures: (a) partial recrystallization in alloy 4, following 1 h anneal at 300 ~ and (b) completerecrystallization in alloy 7, following 10 mm anneal at 525 ~ 90 8O ~...__~~ .~ 70 "--,~ ~ m ~ , o 60 ~ -"~ .......... ........ ~ 6 (AI-Mg) ,D ~- 1(5083) ee 5 0 = 4 (5083+Zr) W 40 w o 7 (5083+Mn) Q c ......... ...... r 9 (5083+Mn+Zr) s_ 30 l .... M Z 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 ~ The recrystallization kinetics at 300 ~ of the various ble for at least 105 s. There is no clear evidence that thealloys are shown in Figure 4, which is directly derived from addition of extra Mn, over and above the commercial baseFigure 3. Alloy 6 recrystallizes completely within 200 sec- alloy (in the absence of Zr), affected the recrystallizationonds at 300 ~ Alloys 1 and 7, with Mn additions, also kinetics within the range evaluated, because alloy 7 con-exhibit complete recrystallization, although with slower ki- taining 1.6Mn exhibited full recrystallization with about thenetics. However, those alloys containing Zr (4 and 9) ex- same kinetics as alloy 1 contains 0.8 pct. Alloy 9 was thehibit a tendency to arrest the full recrystallization at 50 to most resistant to recrystallization at 300 ~ but seems to60 pct, and this arrested recrystallization appears to be sta- recrystallize with characteristics not too different from alloy2950~VOLUME 27A, OCTOBER 1996 METALLURGICAL AND MATERIALS TRANSACTIONS A
  • 5. 1.0 r f 7, a ~ I ! /, a ! e- / / - l ! ,,,,,, 0.8 : A I : tl Q I N r l ! I 9 6 (AI-Mg) m 1 I - J-A-a A 13 1 (5083) m 0.6 I !/ iI ~ / s q o 7 (5083+Mn) O !9 I I /i /" Q & 4 (5083+Zr) L- / Ijl l 9 r 0.4 9 9 (5083+Mn+Zr) O "1 I il ~ /l / ill~ o 10 /9 / d_?t/ t J A ......... 0.2 i / rY" f / B - & ~-- I ~1 I I IIIll 9 l I l s Ill 9 9 9 i 9149 9 9 | |1 9149 0.0 0 0 01 10 2 .10 3 10 4 10 5 Annealing Time (sec)Fig. 4---Recrystallization kinetics, measured at 300 ~ of alloys of different compositions (derived from Fig. 3). 1.0 depending on the initial grain orientation. This possibility //~11 was examined as follows. The average grain size in the as- 08 i received material (12.7-mm thick), measured for alloy 4uA P/" after a 525 ~ h anneal, was 23.4/xm, giving an average / area of 430/xm 2 per grain. While some of the unrecrystal- 06 lized regions in Figure 2(a) have an area approximately equal to the above area, larger areas in excess of 800/zm 2Oe- 0.4 / t/ C // u -" - U I,; 1 have also been observed, indicating the presence of at leastOm iI / two original grains in that region. The former could suggest Io t // that Pa varies from grain to grain, but the latter suggests ; ,i that some other factors also affect recrystallization (viz. in- 0.2 J iu/ / a homogeneity in (fir) and resultant P.). So, while the pos- D sibility of an inhomogeneity in the stored energy cannot be O0 --~ altogether ruled out (at least in the smaller unrecrystallized 10 0 101 10 2 10 3 10 4 10 5 regions), the inhomogeneity in particle densities appears to Annealing time (aec) contribute more strongly.Fig. 5--Recrystalhzatlon kinetics of alloy 9 (5083 + Mn + Zr) at two Alloy 9 does exhibit rapid and complete recrystallizationdifferent temperatures. Rapid recrystallization occurs at 525 ~ without an arrest plateau at 525 ~ as shown in Figure 5, where the 300 ~ recrystallization properties are also shown4, which has a similar Zr content but lower Mn content. for comparison. At this higher temperature, complete re-The partial recrystallization in alloys 4 and 9 suggests a crystallization occurs within about 100 to 500 seconds. Innonuniformity in the particle distribution, with suppression a typical SPF operation at 525 ~ with this alloy, the heat- up and temperature stabilization times are usually at leastof static recrystallization occurring in regions where (fir) 30 minutes. So, it can be concluded that even alloy 9, withexceeds QTr)cr, Therefore, regions with a high local con- t. its extra Mn and Zr additions, would be in a statically re-centration of fine particles, or alternately with very fewlarge particles, could remain unrecrystallized. Figure 2(a), crystallized condition at the start of SPF.which shows regions of fully recrystallized microstructureswith adjoining bands of unrecrystallized material in alloy B. Recrystallized Microstructure4, appears to support the previous statement. It must be noted that local variations in Pd could cause Grain size measurements after different holding times atvariations in (fir),, itself; i.e., in the unrecrystallized 525 ~ are shown in Table II. The term Droc is the grainregions, Pd itself could be lower in addition to P: being size measured in the as-recrystallized condition. It washigher. Upon rolling, Pd might vary from grain to grain, measured after a 10-second anneal for alloy 6, while otherMETALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 27A, OCTOBER 1996~2951
  • 6. 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) Aspect Ratio in DT~ (p~m) c As-Recrystallized D (/~m) D (/zm)Alloy (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 6.54Fig. ~ 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 annealat 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 grow rapidly as well. Grain growth was measured after 16 was found to recrystallize very rapidly and expected to and 24 hours. Figure 6 shows typical optical micrographs2952--VOLUME 27A, OCTOBER 1996 METALLURGICAL AND MATERIALS TRANSACTIONS A
  • 7. 100of the initial as-recrystallized grains in alloys 6 and 9 andof their growth following 1-hour exposures. As seen from the aspect ratios presented in Table II, al-loys 1, 4, 7, and 9, with higher amounts of alloying addi-tions of Mn and Zr, have elongated grains, while alloy 6 ohas an equiaxed grain structure, which is retained as growth goccurs. Grain size results from Table II are graphically rep-resented in Figure 7. The addition of alloying elementsseems to stabilize the grain structure, as seen by the largeamount of grain growth in alloy 6, as opposed to the other D rec ~alloys.C. Particle Sizes and Distribution Electron microscopy studies revealed particles containing 1 ....... ,i ....... .i ........ | ........ | ...... 10 0 101 102 10 3 104 10 5Mn and Cr, which were probably A16(Mn,Cr) and thosecontaining Zr, which appear to be A13Zr. Some of the larger Annealing Time (sac)Mn-bearing particles (>2/zm) also contained Fe. Figure 8 Fig. 7 - - G r a m sizes in different alloys upon anneahng at 525 ~ Alloy 6,shows SEM micrographs of the A16(Mn,Cr) precipitate with no Mn or Zr additions, exhxbits significant grain growth with time.structure in the 5083 base alloy (alloy 1) and 5083 withadditional Mn (alloy 7). The higher Mn addition in alloy 7 different. While an increase in Mn content from 0 (alloy 6)resulted in only a small increase in the density of submi- to 0.8 pct (alloy 1) helps reduce recrystallization kinetics,crometer particles, but the density of large precipitates (>2 it clearly does not lead to its complete suppression. A fur-/zm) increased almost four times, from 0.0017/xm 2 in al- ther increase in the Mn content to 1.6 pct (alloy 7), whileloy 1 to 0.0064/~m -z in alloy 7. producing a slight increase in the number of finer particles, The dispersion of finer particles is shown in the TEM results in a much larger increase in coarser particles. Themicrographs of alloy 1 (5083) and alloy 4 (5083 with Zr) latter effect appears to negate any beneficial effect from thein Figure 9. A13Zr particles present in alloy 4 are much finer former in terms of suppressing static recrystallization.than A16(Mn,Cr) particles, which are present in both alloys.While the A16(Mn,Cr) particles were incoherent, both co-herent and incoherent A13Zr particles were observed. A13Zr B. Particle (f/r) Values and Relation to Grain Sizeparticles were in the form of submicrometer particles, Results of the microstructural analysis clearly indicate awhereas the Mn-containing particles ranged from 0.04 to bimodal distribution of particles (Table III). A simplifying7.1 /zm. Table III shows a compilation of the sizes of the assumption, that the predominant drag pressure is due toparticles, based on the preceding electron microscopy stud- fine particles and that of the larger particles is almost neg-ies and literature, t~6,~7,tsl Volume fractions of the particles ligible, can help make approximate (fir) calculations for thein different alloys, estimated from phase diagrams, are purpose of quantifying Figure 1. This assumption appearsshown in Table IV. These data are utilized in Section IV to be quite reasonable, as reported by Welt and Austin. [~41to estimate a total (fir) for the different alloys. The effect of the large particles, however, cannot always be overlooked, as they do have an indirect influence on IV. DISCUSSION recrystallization by varying the number of potential nucle- ation sites N,. (How significantly N, is varied depends onA. Recrystallization Kinetics the relative density of other potential nucleation sites like The preceding results indicate possible explanations for triple points, shear bands, etc.) Using the volume fractionsthe observations made from Figures 4 through 6. As seen data in Table IV, taking conservative estimates of sizesfrom Figure 8, the addition of excess Mn to 5083 (alloy 7) rA,~(Mn,c,~ = 0.2 /xm and rA~3z,= 0.025 /zm (Table III), andleads to an increase in the number of nucleation sites, due following the Wert and AustinI~"]method of cumulative es-to the fourfold increase in the density of particles >2/xm, timation of (fr), we can get approximate values of (fir) forwhile there is only a small increase in the submicrometer different alloys, as in Table V. The preceding method, inparticle density. The observation that an increase in Mn this case, becomescontent alone (from 0.8 pct in alloy 1 to 1.6 pct in 7) does (f/r)tota, = (f/F)AI6(Mn,Cr) d- 2(f/r)A,,Zr [6]not inhibit recrystallization very much is consistent withthis result (Figure 4). Alloy 9 exhibits a higher resistance The multiplicative factor 2 for the A13Zr particles arisesto recrystallization than 7, because Zr additions are very from earlier work by Nes et al. t~31 showing that coherenteffective in inhibiting recrystallization because of their particles exert twice the drag of incoherent particles. In oursmaller size. Thus, despite being present in lower propor- study, although some large A13Zr particles (with sizestion than Mn, they impose a higher drag. >0.05/zm) were found to be incoherent, for simplicity, we The preceding results, such as the rapid recrystallization treat all Zr-bearing particles as coherent.in alloys 1 and 7, as opposed to 4 and 9, reinforce general The estimated (fir) values, together with the measuredideas of the role of the fine Zr-containing particles in sup- grain sizes from Table V, can be utilized in the graph of Dpression of static recrystallization by restricting the mobility vs (fir) (Figure 10). Grain size measurements indicate aof moving boundaries. The effect of Mn additions is a little very pronounced growth of grains with time, in the case ofMETALLURGICAL AND MATERIALS TRANSACTIONS A VOLUME 27A, OCTOBER 1996~2953
  • 8. 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 (alloy7), 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 possess an initial fine-grained structure, which can be at-the absence of significant growth, in the case of alloys 7, tributed to the high degree of cold work ( - 8 6 pct). But4, and 9. Also shown is a theoretically predicted limiting clearly, in alloy 6 with no Mn and Zr additions (the onlygrain growth size, from Eq. [4], using a of unity. All alloys particles possibly present being those containing the trace2954~VOLUME 27A, OCTOBER 1996 METALLURGICALAND MATERIALS TRANSACTIONS A
  • 9. Table III. Size Distribution of Particles, Obtained from when recrystallized under these conditions. The ratio of the SEM and TEM Studies and Estimates from References 6* density of as-recrystallized grains to the density of particles and 16 through 18"* larger than 2 /zm for the alloy 1 was found to be 22.2, Diameter Diameter while that for alloy 7 with additional Mn was 5.3. This Coarse Fine indicates that there are many more nucleation sites in op- Particles Particles eration, in addition to the large particles in alloy 1, and, to Particle (/Lm) (/zm) a lesser extent, in alloy 7. So, while large particles canA16(Mn,Cr) 1.5 to 7.1 0.04 to 0.2 influence the nucleation of recrystallization, they are not theA13Zr 0.05 to 0.1" 0.01 to 0.025** 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, pos- Table IV. Equilibrium Volume Fractions ( f ) of Various sible that a large particle might nucleate more than one Particles in Modified 5083 Al Alloy at 350 ~ (Estimated grain, but in view of such high ratios, it is quite reasonable from Phase Diagrams, References 19 and 20)* to expect other nucleation sites being operative. A16(Mn,Cr) A13ZrAlloy (Error +__0.0015) (Error _+ 0.0005) C. Estimation of (f/r) .... and Suggested Alloy 1 0.026 -- Modifications 4 0.026 0.0025 From the preceding results, it is apparent that the addition 6 -- -- of Zr and Mn has resulted in an increase in (fir), but this 7 0.052 -- has not been sufficient to prevent the onset of static re- 9 0.052 0.0025 crystallization. In this section, an estimation of the critical *Source of estimatederror: least count of measuring ruler. value of (fir) for subgrain stabilization (at a size of 1 /zm) has been attempted. Table V. Estimation of Total (f/r) for Alloys at 350 ~ (a) One approach toward estimating (fir)on,, similar to that adopted by Nes,tl8j is to consider the equation Alloy Total (fir) 1 0.13 Dgg = a(r/f) [4] 4 0.33 where a is a geometric constant of the order of unity 6 0.00 (=4/3 in Zeners original analysis) and Dgg is now the 7 0.26 9 0.46 size of the subgrain that would be stabilized by the particle dispersion. Therefore, to stabilize a subgrain of 1 /zm, assuming a = 1,elements), the limited particle dispersion cannot stabilize (f/r)cn,- 1 /xm -l [7]this initial fine-grained structure. This is seen in the grainsin this alloy, growing from Dre~ (which is about 9 /zm) to (b) A more rigorous method is to determine the value ofDgg, further growth being restricted by particle dispersion. (fir) at which the driving pressure for recrystallization,The continued grain growth in this alloy between the 1- Pa, equals the Zener drag pressure, ~ , which is theand 24-hour annealing periods could indicate that the equi- desired (fir)r Wert and Austin t141have reported meas-librium Dgg value in this alloy has not been attained after urements of P~ as a function of rolling strain for the 1 hour, due to relatively slow grain mobility. Alternately, A1-0.38 pct Zr system. Using these values as a firstthere could be Ostwald ripening or dissolution of the pre- approximation, for the rolling strain of 1.97 employedcipitates. In contrast to alloy 6, the grain size in alloy 9 is in this study,stable at the level indicated by Dr,r Grain size trends, noted Pd = 0.46 MPa = P_ = 1.5y(f/r)~,t [8]by Vetrano et al.,t6I seem to support this view. It appears, therefore, that with an increase in (fir), there For 3 = 0.3 J/m2, ~131is a transition from grain growth controlled grain sizes (i.e., (f/r)~,, = 1.022 /zm -1 [9]Dgg controlled), as in alloy 6, to one that is controlled bythe as-recrystallized grain size (D J , as in alloy 9. The Assuming a value of (fir)r , of 1 /zm l, the alloy compo-grain growth observed in alloys 1 and 7 is quite small (be- sitions that could lead to this value can be estimated. Theing less than an order of magnitude), which could imply following assumptions have been made.that the particle contents in these alloys are quite close to (a) The alloy has been subjected to an identical TMP, in-the (fir), values for the alloys. A rough sketch of the Dr~o cluding the 350 ~ h preroll aging treatment.profile is shown, based on these considerations. Additional (b) The sizes of the Mn and Zr intermetallic particles aredata points at higher (fir) values could help obtain better assumed to be the same as those used in (fir) estimatesexperimental quantification, but this rough curve is useful reported in Table VI, viz. rA~,tM,,Cr.Fe = 0.2 /xm and ~for illustrative purposes. FAI3Zr = 0.025 tzm. A point of interest is why there is not much change inthe measured as-recrystallized grain size, D~r with (fir). Based on volume fraction data calculated at 350 ~ andThe as-recrystallized sizes in alloys 1, 4, 7, and 9 are almost using the preceding particle sizes, the alloys given in Tablethe same, about 6/xm, indicating a nearly constant density VI have been calculated to have an (flr),o, equal to the cal-of nucleation sites in all of the heavily cold-rolled materials, culated (fir).... of 1 /zm i.M E T A L L U R G I C A L A N D MATERIALS T R A N S A C T I O N S A VOLUME 27A, O C T O B E R 1996--2955
  • 10. 100 ~~= i D rec (sketch) / / [ 1 r 0. L , ~ // .0 o D rec E 6 ~kk /,,/ n D(lh) "wN 10 }" .~.~ g~. ~ =1 ,,," ""~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 at525 ~ (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 graingrowth size, based on Eq. [4], and a rough sketch of the Drr profile. It seems quite unlikely that such fine and uniform particle Table VI. Alloy Concentrations in the Base Alloy Al-dispersions can be achieved in alloys of some of the above 4.75Mg Estimated for a 1-pm Subgrain Stabilizationcompositions, especially those involving the higher Mn Alloy Wt Pct Mn Wt Pct Zrcontents. The results of this study bear out the relativecoarsening of the Mn-containing particle structure, upon in- I 0.00 0.66creasing Mn content (alloy 7 as opposed to alloy I). So, II 0.80 0.58while the success of achieving the desired fine particle dis- III 1.60 0.49 IV 4.20 0.20persion throughout the matrix in these alloys, by the normal v 5.20 0.00ingot casting process, is slight, there is some evidence toindicate the possibility of subgrain stabilization by otherprocessing methods. Higashi et al. 171 have reported a sub-grain structure of about 1 /zm in a A1-5Mg-0.8Zr alloy, onset of recrystallization for the heavily worked materials.formed by PM, in which dynamic recrystallization was ob- While the magnitude of (fir) is helpful, the determinationserved upon SPD at 500 ~ This alloy composition appears of a single value of (fir), based on a single value o f f andquite consistent with alloy I in Table VI. an average value of r, as has been done here, has its limi- tations. The aspect of the microstructure not reflected by aD. Limitations o f Present Approach total (fir) is the particle-size distribution. It is expected that larger particles (greater than 1 to 2/xm), which would be The model developed by Furu et al.[15] can be used to large with respect to the subgrain size, would tend to stim-predict conditions, resulting in complete suppression of ulate recrystallization. The finer particles (less than 1 to 0.5static recrystallization. The authors assume a uniform dis- /~m) may act to inhibit recrystallization by pinning the sub-tribution of particles in the matrix. This is not always valid grain structure. When both large and small particles arein real situations, where a nonhomogeneous distribution of present, the situation is more complex, as there would beparticles may be present. In that case, some regions may competition between these two effects. Regions aroundhave a high enough local value of (fir) to prevent static some large particles may have insufficient fine particles torecrystallization, while others may not, leading to partial prevent nucleation of recrystallization. Conversely, thererecrystallization, as seen in alloys 4 and 9 at 300 ~ (Figure may be regions devoid of large particles, or in which suf-4). ficient fine particle concentrations are present, so as to in- The "Furu graph" (Figures 1 and 10) is helpful in un- hibit nucleation of recrystallization. In the case where bothderstanding results of the study and suggests that the (fir) of these effects may be present, one may observe rapidratio, while increased above that of the 5083 alloy on add- onset of recrystallization but inhibition of full recrystalli-ing Mn and Zr, was not large enough to fully suppress the2956--VOLUME 27A, OCTOBER 1996 METALLURGICAL AND MATERIALS TRANSACTIONS A
  • 11. zation. The results of the research reported here for alloys 2. D.J. Lloyd: Metall. Trans. A, 1980, vol. llA, pp. 1287-94. 3. J.S. Vetrano, C.A. Lavender, C.H. Hamilton, M.T. Smith, and S.M.4 and 9 are consistent with this concept. 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. V. CONCLUSIONS Tokizane, and N. Furushiro, eds., Japan Society for Research on Superplasticity (JSRS), Osaka, Japan, 1991, pp. 447-52. An increase in the Mn content alone from 0.8 to 1.6 pct 5. H. Imamura and N. Ridley: Superplasticity in Advanced Materials,did not impede recrystallization due to the formation of big ICSAM 91, S. Hori, M. Tokizane, and N. Furushiro, eds., Japanprecipitates. The Mn increase, when coupled with the ad- Society for Research on Superplasticity (JSRS), Osaka, Japan, 1991,dition of 0.2 pct Zr, helped suppress static recrystallization, pp. 453-58.at least partially. At higher temperatures, where complete 6. J.S. Vetrano, C.A. Lavender, M.T. Smith, and S.M. Bruemmer: Symp. Proc. of Advances in Hot Deformation Textures and Microstructure,recrystallization occurred, the increase helped stabilize the J.J. Jonas, T.R. Bieler, and K.J. Bowman, eds., TMS, Warrendale,grain structure. An estimate of the critical (fir) ratio re- PA, 1994, pp. 223-34.quired to stabilize a 1-/zm subgrain structure was found to 7. K. Higashi et al.: Superplasticity in Advanced Materials, ICSAM 91,correspond to a significant increase in the desired second- S. Hori, M. Tokizane, and N. Furushiro, eds., Japan Society forphase particle density and a suitable modification in the Research on Superplasticity (JSRS), Osaka, Japan, 1991, pp. 570-80. 8. B.A. Ash and C.H. Hamilton: Scripta Metall., 1988, vol. 22, pp. 277-chemical composition. In practice, it seems unlikely that 82.such fine particle distributions could be achieved in this 9. B.A. Ash and C.H. Hamilton: in Superplasticity and Superplasticalloy system by ingot metallurgy. The model of Furu et Forming, N.E. Paton and C.H. Hamilton, eds., TMS, Warrendale, PA.al.t~sJ was adopted to understand the grain size results ob- 1988, pp. 239-44. 10. E. Nes: Superplastictty, B. Baudelet and M. Suery, eds., Centretained on recrystallization, in terms of (fir) of particles. National de la Recherche Scientifique, Paris, 1985, pp. 7.1-7.14.While their model is a useful qualitative tool, it was found 11. F.J. Humphreys: in Recrystallization "90, T. Chandra, ed., TMS,that the particle size distribution and inhomogeneity in (fir) Warrendale, PA, 1990, pp. 113-22.across the metal sample tends to restrict its applicability. 12. N. Hansen and D.J. Jensen: in Recrystallization 90, T. Chandra, ed., TMS, Warrendale, PA, 1990, pp. 79-88. 13. E. Nes, N. Ryum, and O. Hunderi: Acta Metall., 1985, vol. 33, pp. ACKNOWLEDGMENTS 11-22. 14. J.A. Wert and L.K. Austin: Metall. Trans. A, 1988, vol. 19A, pp. 617- This research was supported by a grant from Pacific 25. 15. T. Furu, K. Marthinsen, and E. Nes: Mater. Sci. Forum, 1993, vols.Northwest National Laboratory. The work of JSV was sup- 113-115, pp. 41-52.ported by the Materials Science Division, Office of Basic 16. T.R. McNelley and S.J. Hales: Superplasticity in Aerospace 11, T.R.Energy Sciences, United States Department of Energy McNelley and C.H. Heikkenen, eds., TMS, Warrendale, PA, 1990,(DOE), under Contract No. DE-AC06-76RLO 1830. The pp. 207-22.authors would also like to thank Curt Lavender and Mark 17. T.R. McNelley and P.N. Kalu: Superplasticity in Advanced Materials, ICSAM 91, S. Hori, M. Tokizane, and N. Furushiro, eds., JapanSmith, Pacific Northwest National Laboratory, for helpful Society for Research on Superplasticity (JSRS), Osaka, Japan, 1991,discussions. 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 REFERENCES and Co., London, vol. 8, pp. 258-65 and 342-45. 20. L.F. Mondolfo: Aluminum Alloys--Structure and Properties, 1. B.J. Dunwoody, R.J. Stacey, and A.J. Barnes: Superplasticity in Butterworth and Co., London, 1976, pp. 250, 324, 338, and 413. Metals, Ceramics and Intermetallics, Materials Research Society, 21. E.S. Puchi and G. Rojas: Mater. Sci. Forum, 1993, vols. 113-115, Pittsburgh, PA, 1990, pp. 161-66. pp. 509-14.METALLURGICAL AND MATERIALSTRANSACTIONS A VOLUME 27A, OCTOBER 199~-2957