THE SECONDARY HARDENING PHENOMENON IN

STRAIN-HARDENED MP35N ALLOY

S. ASGARI{, E. EL-DANAF, E. SHAJI, S. R. KALIDINDI and R. D. DOHERTY{Department of Materials Engineering, Drexel University, Philadelphia, PA 19104, U.S.A.

(Received 18 November 1997; accepted 19 June 1998)

AbstractÐMechanical testing and microscopy techniques were used to investigate the in¯uence of aging onthe structure and strengthening of MP35N alloy. It was con®rmed that aging the deformed material at6008C for 4 h provided additional strengthening, here referred to as ``secondary hardening'', in addition tothe primary strain hardening. The secondary hardening phenomenon was shown to be distinctly di�erentfrom typical age hardening processes in that it only occurred in material deformed beyond a certain coldwork level. At moderate strains, aging caused a shift in the entire stress±strain curve of the annealed ma-terial to higher stresses while at high strains, it produced shear localization and limited work softening.The secondary hardening increment was also found to be grain size dependent. The magnitude of the sec-ondary hardening appeared to be controlled by the ¯ow stress in the strain hardened material. A model isproposed to explain the observations and is supported by direct experimental evidence. The model is basedon formation of h.c.p. nuclei through the Suzuki mechanism, that is segregation of solute atoms to stack-ing faults, on aging the strain hardened material. The h.c.p. precipitates appear to thicken only in the pre-sence of high dislocation density produced by prior cold work. # 1998 Acta Metallurgica Inc. Published byElsevier Science Ltd. All rights reserved.

1. INTRODUCTION

MP35N (35Co±35Ni±20Cr±10Mo) is a cobalt±nickel base superalloy extensively used for low tomoderate temperature applications (up to 5008C)requiring high strength, high toughness, and highcorrosion resistance. When fully annealed at tem-

peratures above 8508C, MP35N is a single phasesolid solution with f.c.c. (face-centered cubic) struc-

ture. The alloy shows an impressive strain harden-ing response at room temperature and is

commercially strengthened mainly by cold working.A detailed investigation into the mechanisms of

hardening at room temperature deformation of anumber of low stacking fault energy (SFE) alloys,

including MP35N, was recently published [1]. Itwas shown that the strain hardening behavior of

MP35N originated from extensive formation of de-formation twins and the ability of twin boundariesin hindering crystallographic slip. The alloy can be

further strengthened by heat treating the deformedmaterial. The commercial practice is to heat treat

the 48±53% cold drawn MP35N at 5938C for 4 hfollowed by air cooling (see, e.g. Ref. [2]). This

causes a remarkable 40% increase in yield strength,from 1385 to 1780 MPa, while maintaining a high

ductility of 50% reduction of area in a round bartensile test. In the present paper, the hardening

e�ect caused by post-deformation aging of MP35N

alloy is referred to as ``secondary hardening'' to dis-

tinguish it from the ``primary hardening'' achieved

by room temperature deformation through strain

hardening. Aspects of the secondary hardening

phenomenon have been investigated in previous

studies [2]. It has been reported that a material

deformed to higher cold work levels shows larger

magnitudes of secondary hardening after aging [3, 4].

Also, it was found that heat treating the as-

annealed (or as-quenched) alloy at 5938C even for

very long times (beyond 100 h) produced no

hardening [3]. This behavior distinguishes the sec-

ondary hardening phenomenon from the usual pre-

cipitation hardening processes. In an ordinary

precipitation hardening process, deformation is not

required for precipitation to occur, although

deforming a supersaturated alloy can change the

precipitation kinetics and the magnitude of the

resulting strengthening. As explained later in this

paper, the structural changes occurring during

aging of deformed MP35N alloy are distinctly

di�erent from those of typical age hardening reac-

tions. For this reason, we will use the term ``aging''

in this paper only to describe the thermal treatment,

i.e. isothermal heating at 5938C for 4 h. Also, the

term ``secondary hardening'' is substituted for age

hardening throughout the paper to distinguish

between the process in MP35N alloy and conven-

tional age hardening.

The microstructural origin of the secondary hard-

ening phenomenon, has been investigated by several

authors. Graham [5] originally proposed two

Acta mater. Vol. 46, No. 16, pp. 5795±5806, 1998# 1998 Acta Metallurgica Inc.

Published by Elsevier Science Ltd. All rights reservedPrinted in Great Britain

1359-6454/98 $19.00+0.00PII: S1359-6454(98)00235-3

{Present address: Department of Materials Engineering,University of Tehran, Tehran, Iran.

{To whom all correspondence should be addressed.

5795

alternative mechanisms: (1) precipitation of Co3Mo

(ordered h.c.p.) during aging, and (2) solute parti-

tioning between the f.c.c. matrix and the strain-

induced martensite phase believed to form during

room temperature deformation. The partitioning

model was further developed by Singh and

Doherty [2] who proposed that upon aging, parti-

tioning occurred locally adjacent to the pre-existing

martensitic h.c.p. plates causing an increase in the

stress required for nucleating new plates. No expla-

nation was o�ered, however, for the magnitude of

the secondary hardening produced by this process.

TEM studies [4, 6] did not con®rm the presence of

the Co3Mo precipitates originally proposed by

Graham.

Moreover, detailed TEM investigations [1, 6] have

shown that the plate-like features formed during

room temperature deformation of MP35N alloy are

only deformation twins and are not strain-induced

martensite. This is in agreement with previous TEM

observations reported by Raghavan et al. [7] on

commercially drawn MP35N. In addition, recent

work [6] on the secondary hardening e�ect in ultra-

®ne MP35N alloy (1 mm grain size) showed that

while deformed samples of ®ne grain size alloy con-

tained no strain induced plate-like features (twins

or martensite) up to a strain of about 0.3 in simple

compression, aging the material deformed to this

strain level still produced signi®cant secondary

hardening. The lack of martensite plates in the

deformed MP35N samples, rules out solute parti-

tioning to pre-existing martensite plates as a poss-

ible mechanism for secondary hardening. Raghavan

et al. [7] investigated the structure of the commer-

cially drawn MP35N alloy (59% cold drawn) both

in unaged and aged conditions. These authors con-

cluded that the precipitation of an h.c.p. phase

formed during aging was the main cause of the sec-

ondary hardening phenomenon. The TEM data pre-

sented in support of this idea [7], however, was not

conclusive for the following reasons:

1. The [110] di�raction pattern presented in Fig. 5

of Ref. [7] showed no intensity maxima at the

expected locations for an h.c.p. phase, although

strong streaking along the [111] direction was

seen in the di�raction pattern (this streaking

may have been caused by a large density of

stacking faults and microtwins).

2. The ``h.c.p. spot'' chosen for dark ®eld imaging

was not symmetric with respect to the 000 spot

and showed signi®cant intensity variation (such

non-symmetric intensity distribution may occur

due to slight beam deviations from the zone axis

or as a result of aperture displacement).

In addition, no explanation has been o�ered to

account for the role of deformation prior to aging

which is now found to be an essential ingredient for

the secondary hardening phenomenon.

Not only is there no understanding of the micro-structural origin of the secondary hardening, but a

detailed characterization of the secondary hardeningphenomenon is also lacking, at least in the scienti®cliterature. In particular, no systematic study on the

in¯uence of cold work level and grain size e�ects onthe magnitude of secondary hardening has beenpublished. In addition, the e�ect of the secondary

hardening phenomenon at high strain levels onpost-aging deformation behavior of MP35N alloy isnot well documented. A clearer understanding and

characterization of the secondary hardeningphenomenon in MP35N alloy seems to be an essen-tial step for improving deformation processing andexploiting the full potential of this interesting alloy

and other related alloys [2, 4, 6]. Finally, the lack ofan explanation for the remarkable strengtheningcaused by aging is a challenge to physical metal-

lurgy.In the present investigation, mechanical testing

and microstructural studies were employed to

characterize the mechanical behavior of coldworked and aged MP35N and to obtain a morecomprehensive understanding of the physical origin

of the secondary hardening phenomenon. Thespeci®c goals of this investigation were (i) tocharacterize the e�ects of di�erent amounts of coldwork on the magnitude of secondary hardening in

1, 35, and 300 mm grain size MP35N, and (ii) tostudy the microstructural origin of the secondaryhardening e�ect using samples deformed to smaller

strains compared to those used in previous studies,e.g. Refs [2, 5, 7]. Use of more lightly deformedsamples should minimize the complications arising

from heavily deformed structures investigated inprevious studies and increase the chance of bettercharacterization of the structural changes caused byheat treating the deformed alloy.

2. EXPERIMENTAL METHODS

Commercially drawn rods of MP35N alloy pro-vided by SPS Technologies, Inc., Jenkintown,

Pennsylvania, were used to produce annealed ma-terial with average grain sizes of 300, 35, and 1 mm.The 300 mm material was produced by heat treatingthe as-received material at 12008C for 5 h and the

35 mm samples were prepared by heat treating at10008C for 1 h. Details of the process used to pro-duce the 1 mm grain size material are described in

detail elsewhere [2, 6]. In outline, multiple cycles ofmoderate deformation and recrystallization at8008C were used. Cylindrical samples with diam-

eters of about 5±8 mm, and lengths of about 7±12 mm were tested in simple compression using acomputer-controlled servo-hydraulic MTS mechan-

ical testing unit. At room temperature, a constanttrue strain rate of ÿ0.001/s was maintainedthroughout the tests. The specimen ends were lubri-cated using Te¯on sheets and high pressure grease

ASGARI et al.: SECONDARY HARDENING5796

to minimize frictional e�ects. As in the earlier

study [1], by frequent relubrication at small incre-

ments of strain, it was possible to minimize the fric-tional e�ects as indicated by the absence of

``barreling'' and by the high level of reproducibility

of the stress±strain data. True stress±true strain

curves were computed from raw data acquired in

terms of load and displacement, after being cor-rected for the compliance of the testing machine [8].

The change in yield strength caused by aging was

measured by comparing the yield stress of the aged

sample with the ¯ow stress prior to aging. In somecases, particularly at high strains, larger samples

were ®rst deformed to a desired strain using the

mechanical testing facilities at SPS Technologies

Inc. Identical compression samples of smaller diam-

eter were then cut from these samples using an elec-tric discharge machine (EDM) for further

compression testing in our laboratory. Samples for

optical microscopy were sectioned parallel to the

loading direction, mechanically ground, etched in asolution of FeCl3 in HNO3, and examined in an

Olympus PMG3 microscope. Thin foils for TEM

studies were prepared by standard electropolishing

using a Fischone twin-jet unit. A solution of 60 mlH2SO4, 15 ml H3PO4, and 240 ml methanol was

used at a temperature of 08C and an operating vol-

tage of 10 V. A Philips 400T STEM operating at

120 kV and a JEOL 100CX II TEM operating at

100 kV were used for electron microscopy studies.

3. RESULTS

3.1. Secondary hardening in simple compression tests

3.1.1. Commercially drawn material. Initial testson MP35N commercially cold drawn to anelongation of 48% gave a large increment of 0.2%

proof stress, from 1250 to 1750 MPa after aging thecold drawn sample at 5908C for 4 h. Very similarresults have been previously reported for tensiletests on similar material [2].

3.1.2. Initially compressed material. The main stu-dies to characterize secondary hardening in thisstudy were compression tests performed on recrys-

tallized material deformed by uniaxial compressionto di�erent prior strains followed by aging andfurther uniaxial compression. Figure 1 shows

examples of the experiments on alloys recrystallizedto a grain size of 35 mm. The curve labeled I is thestress±strain behavior of annealed MP35N alloy in

simple compression with no aging [1]. Curve II isthe stress±strain response of a sample which, afterbeing deformed in compression to a true strain of0.28, was aged and further deformed in simple com-

pression. This gave a small increment in ¯ow stressseen but with the subsequent true stress±straincurve parallel to that of the unaged material. After

about 13% deformation, the compression test wasinterrupted and the sample was aged again. Thecurve labeled III in Fig. 2 shows the continued

stress±strain behavior of the sample, following the

Fig. 1. True stress±true strain response of MP35N alloy with 35 mm grain size deformed in simple com-pression: (I) fully annealed; (II) prestrained to 0.28 and aged; (III) response of the same sample used in

(II) after reaging and resuming the test; (IV) prestrained to 0.58 and aged.

ASGARI et al.: SECONDARY HARDENING 5797

second aging. It had an additional increment of¯ow stress and again showed similar strain harden-

ing rates as the unaged material of curve I. This

behavior was always found for samples moderately

deformed before aging. Material deformed to higher

prestrains before aging showed a very di�erentbehavior however. This is illustrated in curve IV for

MP35N prestrained without aging to a strain of

0.56, at which point it was aged and further com-

pressed. Curve IV shows a larger increment of ¯owstress, continuing the trend of a larger increment of

secondary hardening with increased prior strain

hardening, but now there was a limited period of

strain softening, typically lasting for a strain incre-

ment of about 0.1 before further strain hardeningcontinued. This short period of strain softening was

also found for the commercially drawn material

when studied by compression [3]. Prior studies

using tensile testing of drawn material [2] did not

see this behavior since strain softening in a tensiletest will lead to local necking. Once necking occurs

further observation of the true stress±true strain re-

lationship cannot be measured from load/displace-

ment measurements. Detailed studies of thise�ect [3] have shown that these local strain soften-

ing e�ects occur with strain localization in shear

bands. After such periods of local softening associ-

ated with strain localization the sample scale stress±

strain curves even in compression no longer providematerial data since the strain is no longer homo-

geneous. Data up to the maxima appear, however,

to come from homogeneously deforming material.

Figure 2 summarizes the data obtained for the

secondary hardening increments for a range ofstrains in the alloys with the three di�erent grain

sizes, 1, 35, and 300 mm. The increase in stress, Ds,between the ¯ow stress of the material deformed to

a given strain, before aging and the yield stress

after aging, is plotted in Fig. 2. The data shown are

for samples aged only once, as in curves II and IV

of Fig. 1. No data for multiply aged samples are

shown. Also marked, by vertical arrows, are the

samples that ®rst showed the strain softening seen

in curve IV in Fig. 1. The strains at which strain

softening was ®rst seen after aging were 0.5 for

1 mm grain size, 0.6 for 35 mm grain size, and 0.9

for 300 mm grain size. In each case, however, the

®rst onset of strain softening occurred in samples

with a prior ¯ow stress of 1600±1700 MPa and a

secondary hardening increment of about 200 MPa

as seen in Fig. 2.

Figure 2 also shows one of the most signi®cant

results of this study. The magnitude of the second-

ary hardening appears to be controlled mainly by

the prior ¯ow stress. Samples deformed to ¯ow

stresses of less than 700±900 MPa showed no sec-

ondary hardening on aging. This result has been

seen qualitatively previously [3] where fully

annealed samples aged at 5908C for times up to

300 h showed no measurable increase in microhard-

ness with aging. In addition, tensile tests of samples,

with 35 mm grain size, pulled to small strains

(e< 0.2) also showed no measurable secondary

hardening [3]. Secondary hardening is seen in Fig. 2

to show a nearly linear increase with prior ¯ow

stress in compression, from Ds of near zero at a

prior ¯ow stress of about 1000 MPa to a saturation

value of Ds of about 300 MPa at a prior ¯ow stress

of about 1700 MPa.

Fig. 2. Variation of the increase in yield strength, Ds, of MP35N alloy due to aging at 5938C for 4 hwith ¯ow stress prior to aging for the three di�erent grain sizes, 1, 35, and 300 mm.

ASGARI et al.: SECONDARY HARDENING5798

3.2. Microscopy

To study the microstructural origin of the sec-ondary hardening phenomenon, samples of 35 mmgrain size MP35N deformed in the strain range of0.09 (¯ow stress of 750 MPa) to 0.25 (1100 MPa) in

simple compression both with and without agingwere used. This strain range was chosen for TEMstudies since the secondary hardening e�ect gradu-

ally developed with increasing strain and prior ¯owstress in this range. Furthermore, at these low

strains, complications arising from the heavily

deformed structure of the commercially drawn ma-

terial used by previous investigators [2] could be

minimized and the structural changes caused by

aging can be examined more clearly. Figures 3(a)

and (b) show the dislocation structures of two

samples deformed to strains of 0.09 and 0.15, re-

spectively. These structures are typical of deformed

low SFE materials characterized by planar arrays of

dislocations, in contrast to cell formation character-

istic of medium to high SFE materials [9]. TEM

micrographs shown in Figs 4(a)±(c) are high magni-

Fig. 3. TEM micrographs of annealed MP35N alloy of 35 mm grain size deformed in simple com-pression to a strain of (a) 0.09 or (b) 0.15.

ASGARI et al.: SECONDARY HARDENING 5799

®cation images obtained from samples of MP35Nalloy deformed to strains of 0.09, 0.15, and 0.25,

and then aged at 5938C for 4 h. It is important tonote that the ®rst two samples showed no measur-

able hardening while the sample deformed to 0.25,showed a small but ®nite hardening of about

20 MPa. TEM observations in these deformed andaged samples did not show, as previously noted [2],

Fig. 4. TEM micrographs of 35 mm grain size MP35N deformed in simple compression to a strain of(a) 0.05, (b) 0.15, or (c) 0.25, and then aged.

ASGARI et al.: SECONDARY HARDENING5800

any sign of a general precipitation. The striking fea-ture observed in these structures seen in Figs 3 and

4 was the presence of extended stacking faults.These faults are readily identi®ed by their character-istic fringe patterns. The density of these stacking

faults increases with increasing the plastic strainprior to aging. The next important observation wasthe clear increase of the widths of the stacking faults

on aging. In the lightly compressed samples (0.05and 0.15) that showed no measurable hardening thelength increase was limited. With the sample that

showed a small hardening e�ect [Fig. 4(c)], thestacking faults grew much larger and showed exten-sive overlapping on apparently closely spaced paral-lel slip planes.

4. DISCUSSION

A unique property of the MP35N alloy, which isof major commercial importance, is the signi®cant

increase in the yield stress achieved by aging thecold worked material at about 6008C. The resultsreported here, combined with the earlier obser-vations described in Section 1, show that the sec-

ondary hardening of this alloy is very di�erent fromconventional precipitation hardening. In particular,Fig. 2 shows that the magnitude of the secondary

hardening e�ect is clearly dependent on the prior¯ow stress level.Variation of Ds with ¯ow stress is well rep-

resented by a multi-stage response (Fig. 2). In the®rst stage, no measurable hardening is obtained byaging the deformed material. During the second

stage of secondary hardening the magnitude of Dsincreases slowly with ¯ow stress. This is followedby stage three, a rather rapid almost linear increasein the magnitude of Ds at higher ¯ow stresses. The

fourth and ®nal stage, saturation, characterized bya constant Ds, then occurs.The TEM micrographs shown in Figs 3 and 4

provide important clues to the physical origin ofthe secondary hardening phenomenon. The strikingfeature seen in these micrographs is the presence of

a high density of extended stacking faults in theaged material [see Figs 4(a)±(c)]. It is important tonote that while extended faults are observed in allthree samples, only the sample in which a high den-

sity of overlapping faults occurs shows a measur-able hardening [Fig. 4(c)]. One mechanism by whichstacking faults in concentrated f.c.c. alloys can

extend as a result of heat treatment is the Suzukimechanism [10, 11]. This process is driven by thedecrease in the stacking fault energy by solute seg-

regation and has been reported in a number ofalloy systems [12±16]. Since a stacking fault in f.c.c.structures is a thin h.c.p. layer, the solute content in

this h.c.p. fault should be di�erent from that in thef.c.c. matrix. This is particularly true for thosesolute atoms which show preferential solubility forh.c.p. structure (such as Cr and Mo in MP35N

alloy). As a result, a driving force should exist for

solute atoms to segregate either to, or away from,stacking faults in f.c.c. structures. This e�ect maybe signi®cant in low SFE metals such as MP35N in

which dissociation of perfect dislocations to partialdislocations bounding a faulted region occurs easily.

This is shown schematically in Fig. 5. Figure 5(a)corresponds to a moderate SFE metal where separ-ation of Shockley partials is small. In low SFE

metals, however, the partials are widely separated.As a result the cross-slip process of screw dislo-

cations becomes di�cult leading to the well-knownplanar arrangement of dislocations in low SFE

metals, clearly seen in Figs 3(a) and (b). When theSuzuki mechanism occurs, the separation of partialsshould become wider due to the fall in stacking

fault energy absorbing h.c.p. stabilizing elements(e.g. Mo and Cr in the case of MP35N) at the fault.

In general, the Gibbs adsorption isotherm [17]predicts a decrease in interfacial energy, dgSFE, withincreased activity, aB, of any solute that shows posi-

tive adsorption, i.e. GB>0, at an interface:

dg � dgSFE � ÿGBRT dln aB: �1�That is, any solute that preferentially adsorbs at thestacking fault will lower the stacking fault energy

and increase the separation of the partials. In ad-dition, other solutes such as nickel, that stabilizethe f.c.c. structure may segregate away from the

faulted region, GB<0 and DaB<0, which will alsolower gSFE. The stacking fault energy of the isolated

faults seen in Figs 4(a) and (b) appears to remainpositive since the increase in the size of the stackingfaults is limited.

Figure 4(c) shows a di�erent e�ect. In this ®gurewhich corresponds to a higher ¯ow stress (dislo-

cation density) prior to aging, some of the stackingfaults appear to grow almost without limitÐat least

in regions where they overlap other stacking faults.This implies that overlapping stacking faults canform without any increase of chemical free energy.

In other words, in the presence of a high dislocationdensity, chemically stable stacking fault energy

becomes negative and so thin hexagonal regionsseem to nucleate in the f.c.c. matrix. It is this

change that correlates with the small increment of

Fig. 5. Schematic presentation of separation of theShockley partials: (a) medium SFE metal; (b) lowerSFE metal; (c) lower SFE metal assisted by the Suzuki

mechanism.

ASGARI et al.: SECONDARY HARDENING 5801

secondary hardening observed. Previous studies on

the role of stacking faults in f.c.c. twinning andf.c.c. 4 h.c.p. transformation also suggest that dis-location density is an important factor in these

processes [18±20]. A twin can be produced in anf.c.c. crystal by passage of Shockley partials onevery {111} plane of the matrix, while passage of

such partials on every other {111} plane creates anh.c.p. phase [21]. Formation of an f.c.c. twin or

h.c.p. phase of macroscopic size may be consideredas a nucleation and growth process. Once a twin orh.c.p. embryo is formed by dislocation interactions,

it may then thicken by overlapping of other nucleiformed on the {111} planes parallel to its habitplane.

Based on the observations reported here, we pro-pose the following model for the secondary harden-

ing phenomenon in MP35N alloy. On aging thestrain hardened alloy, isolated dislocations undergothe Suzuki mechanism and become more extended

by absorbing Mo and Cr atoms, known to stabilizethe h.c.p. phase [22, 23] and by desorbing Ni atoms.At low to moderate dislocation densities, these two

layer h.c.p. regions do not appear to be capable offurther lengthening and thus cannot be considered

as true precipitates of a stable second phase. So,these regions behave essentially solely as stackingfaults and do not o�er any measurable hardening

[see Figs 4(a) and (b)]. However, at higher dislo-cation densities, where stacking faults may be onlyone or two {111} planes apart, potential h.c.p.

nuclei appear to form and grow by Suzuki segre-gation and produce stable h.c.p. platelets which canapparently lengthen without limit if they overlap.

This new model of secondary hardening wastested with two further experiments. In the ®rst test,

a 35 mm grain size sample was deformed to a strainof 0.48 in simple compression and the sample wassubsequently aged and thin foils were prepared for

TEM investigation. This sample had a strain har-dened ¯ow stress of about 1500 MPa and a second-ary hardening increment of about 150 MPa (see

Fig. 2). Figures 6(a) and (b) are a bright ®eld (BF)image and a [110] di�raction pattern (DP), respect-

ively, and Fig. 6(c) gives the indexed pattern. Theshort very thin plates (only about 5 nm thick)observed in the BF image are distinctly di�erent

from the long, thick, twin plates (marked by T)formed during room temperature deformation [1].The thin plates have an h.c.p. structure and are

formed on two intersecting {111} planes withcharacteristic orientation relationship with the f.c.c.

matrix, as revealed by the selected area di�ractionpattern shown in Fig. 6(b). The di�raction patternwas taken from a twin free region. The precipitates

seen in Fig. 6, though very thin (a few nm), areclearly thicker than a stacking fault (two {111}planes seen edge on). As discussed in a recent

paper [1], a detailed TEM study on annealedMP35N samples deformed to similar (and higher)

strain levels but, without any subsequent aging did

not show any evidence of an h.c.p. phase. We,therefore, believe that the ®ne h.c.p. plates observedin Fig. 6 are formed during the aging process. This

result strongly suggests that formation of a signi®-cant density of stable h.c.p. plates in aged MP35Nalloy after aging requires a very high dislocation

density in the strain hardened material prior toaging. It would be desirable to demonstrate solute

enrichment in the h.c.p. platelets seen in Fig. 6.These plates, however, appear to be too thin toallow reliable use of available analytical TEM tech-

niquesÐparticularly if the solute changes are small.The second test was also performed on a sample,

showing a secondary hardening increment of

200 MPaÐwhich is less than the saturationstrengthening of 300 MPa in compression. This

sample is very similar to the sample studied inFig. 6. After aging for 4 h at 5938C to give the sec-ondary hardening of about 200 MPa, additional

aging was applied for up to 100 h at the same tem-perature. This additional aging time did not lead toany further hardeningÐthat is the ¯ow stress of the

sample was una�ected by this extended aging. Thisresistance to extra hardening occurred in spite of

the potential of the alloy for additional hardeningup to 300 MPa, corresponding to the materialdeformed to higher ¯ow stresses prior to aging (see

Fig. 2). This result indicates that the magnitude ofsecondary hardening is controlled by the prior ¯owstress (the dislocation density) and is not a�ected by

additional aging time. This behavior is in clear con-trast to typical di�usion-controlled precipitationhardening where the rates of hardening can be

a�ected by changing the dislocations that oftennucleate the strengthening precipitates [21] but

supersaturated solid solution will on aging continueto form precipitates. In the present alloy, the beha-vior is di�erent. If, the secondary hardening arises

from h.c.p. precipitation, the amount of precipitationthat can occur in this alloy is determined by theprior dislocation density and it appears to be inde-

pendent of increased aging time, even though thealloy remains supersaturated with h.c.p. forming

solute. The evidence for supersaturation with a sec-ondary hardening of only 200 MPa is supported bythe ability if deformed more, in compression to

show secondary hardening of up to 300 MPa.Growth kinetics of hexagonal precipitates in an

f.c.c. matrix have been studied in detail by several

investigators as reviewed by Doherty [21]. It isknown that the thickening of coherent h.c.p. pre-

cipitates is controlled by the formation and mi-gration of growth ledges [24, 25]. These areShockley partial dislocations. These partial dislo-

cations appear to be generated in the Al±Ag alloysystem by impingement of precipitates on di�erentmatrix habit planes, see Rajab and Doherty [26, 27].

The nucleation mechanism for forming the neededgrowth dislocations, however, remains unclear [27].

ASGARI et al.: SECONDARY HARDENING5802

Without further experimental study of the present

system, more detailed discussion would be largely

speculation, although the present results do appear

to show that the only way thickening of the h.c.p.

precipitates in MP35N can occur is by the presence

of a high density of pre-existing dislocations pro-

duced during pre-aging deformation. These dislo-

cations appear to provide the ledges required for

thickening of the h.c.p. precipitates in the f.c.c.

matrix [24] during the aging process. This, we

believe, may explain the rapid increase in the mag-

nitude of Ds with ¯ow stress during the second

stage of secondary hardening (Fig. 2). Figure 6(a)

suggests that the lengthening of the precipitates is

Fig. 6. TEM results obtained from annealed MP35N deformed in simple compression to a strain of0.48 and subsequently aged: (a) bright ®eld image; (b) di�raction pattern; (c) the indexed pattern.

ASGARI et al.: SECONDARY HARDENING 5803

inhibited by the impact of the thin plates on the

broad faces of precipitates forming the so-called

``T'' junctions [25]. On a coarser scale, lengthening

of the h.c.p. precipitates can also be halted by the T

junctions with the deformation twins on nonparallel

{111} planes.

Fig. 7. E�ect of multiple aging on the true stress±true strain response in simple compression ofannealed MP35N alloy (35 mm grain size): (a) prestrained to 0.62; (b) prestrained to 1.04.

ASGARI et al.: SECONDARY HARDENING5804

A further signi®cant feature of the secondary

hardening increment is its saturation at about

300 MPa. This saturation occurs, in compression

deformed samples, at a prior ¯ow stress of about

1800±2000 MPa (Fig. 2). The precipitation model

for secondary hardening predicts that saturation,

limited by the solute availability, will occur since

only a ``lever rule'' amount of second phase can

form. So the saturation seen during the ®nal stage

of secondary hardening of Fig. 2 is a necessary fea-

ture of the present model. To evaluate this aspect

of secondary hardening, two samples with a grain

size of 35 mm were strained in simple compression

to 0.62 and 1.04, respectively. Note that the sample

compressed to a strain of 0.6 has a prior ¯ow stress

of 1600 MPa, less than the saturation limits of

Fig. 2, while the sample deformed to a strain of

1.04 is well beyond the saturation limit. Figures

7(a) and (b) show the e�ect of multiple aging on

the true stress±true strain response of these samples.

In each ®gure, the true stress±true strain response

of the material deformed without aging is also

plotted as a reference. On aging after a strain of

0.6, there was a yield stress increase of about

200 MPa which, as expected, was less than the sat-

uration level of 300 MPa. Based on the precipi-

tation model of the secondary hardening, this

condition requires that the alloy would still be

supersaturated. After aging, the sample was again

compressed by 10%, to a total strain of 0.72

[Fig. 7(a)] and then reaged and retested. As

expected for a still supersaturated alloy, a small ad-

ditional hardening response was found [Fig. 7(a)].

Further compression by an additional 8% (a total

strain of 0.80) followed by aging, however, resulted

in no further secondary hardeningÐmerely a small

strain aging e�ect was seen [Fig. 7(a)]. Such strain

aging e�ects were previously reported [3] in samples

lightly deformed in tension. Figure 7(b) shows the

results of multiple aging tests applied to the sample

prestrained to 1.04. As expected, the ®rst aging

causes a larger increase in yield stress of the sample

prestrained to 1.04 (Ds of about 300 MPaÐthe sat-

uration value) compared to that of the sample pre-

strained to 0.62 with a Ds of approximately

200 MPa [Fig. 7(a)]. Further aging of this second

sample, however, did not cause any additional

hardening [curve III of Fig. 7(b)] except for the

small strain aging e�ect always seen in saturated

and reaged samples. A further reaging test, curve

IV, showed the same result. These results, we

believe, strongly support the solute depletion hy-

pothesis for the saturation e�ect during the ®nal

stage of the secondary hardening phenomenon.

The results presented here reveal some important

aspects of the secondary hardening e�ect in MP35N

alloy, and the proposed model appears to be in

agreement with most of the experimental data.

However, at least two features of the secondary

hardening process remain somewhat unclear. Theseare:

1. The limited amount of work softening and theassociated strain localization (shear banding)seen during post-aging deformation at high cold

work levels seen in Figs 1 and 7. Such work soft-ening e�ects, however, did not occur in theunaged samples deformed to su�ciently high

stresses to produce shear bands [1]. This strainsoftening may be related to dislocation cutting ofthe very thin precipitates found in stage 3 of sec-

ondary hardening. Gutaez and Penisson [28]have described the details of the dislocationdamage seen in thin h.c.p. precipitates in an

f.c.c. Ni±Co±Cr±Nb alloy during small strains.2. The higher magnitude of secondary hardening,

about 500 MPa, seen in commercially drawnsamples [2] compared to the saturation level of

only 300 MPa seen in the uniaxially compressedsamples.

5. CONCLUSIONS

1. Strengthening of MP35N alloy by aging, second-ary hardening, requires signi®cant prior strain

hardening. At moderate strains, aging causes ashift in the entire stress±strain curve of theannealed material to higher stresses and, at

higher strains, it promotes shear banding duringsubsequent deformation.

2. At a given grain size, no measurable secondary

hardening occurs below a certain plastic strainprior to aging. For prior strains larger than thislimit, the magnitude of the secondary hardeningphenomenon, Ds, increases with plastic strain

and reaches saturation at high strains. At agiven plastic strain, the magnitude of the second-ary hardening e�ect decreases with the increase

in grain size, i.e. the Ds±e curve is shifted tohigher strains. Flow stress prior to aging seemsto be the controlling parameter in the secondary

hardening phenomenon. For the wide range ofgrain sizes used in this study, variation of themagnitude of the secondary hardening e�ect, Ds,with ¯ow stress prior to aging is well represented

by a four stage response. During the ®rst stageno measurable hardening occurs in the deformedmaterial. Stage 2 is characterized by a small

increase in Ds with ¯ow stress, while stage 3shows a rapid and nearly linear increase in Dswith ¯ow stress. In the ®nal stage the magnitude

of Ds saturates.3. Based on the observations, a model for the sec-

ondary hardening phenomenon is proposed and

is supported by experimental evidence. It issuggested that secondary hardening arises instage 2, from solute segregation leading to nega-tive values of the stacking fault energy by Suzuki

ASGARI et al.: SECONDARY HARDENING 5805

solute segregation to the stacking faults. TheSuzuki segregation only causes any measurable

strengthening when there is a high enough den-sity of dislocations for the stacking faults to besu�ciently close to overlap. The major harden-

ing, stage 3, arises from the formation of verythin precipitates of a hexagonal phase. Unlikeother precipitation hardening reactions, including

the formation of hexagonal Ag2Al in an f.c.c.matrix in Al±Ag alloys, the present precipitationreaction appears to require a very high density

of dislocations. These are needed both fornucleation of the hexagonal nuclei and also forsubsequent thickening of these nuclei into stableh.c.p. precipitates of a ®nite thickness.

4. The saturation of secondary hardening at highprior ¯ow stresses is attributed to the depletionof the matrix from h.c.p. stabilizing elements

during the subsequent aging process.5. With increments of secondary hardening exceed-

ing about 200 MPa a limited amount of strain

softening is seen leading to extensive shear band-ing. It is suggested that this may be related todislocation cutting of the very thin h.c.p. precipi-

tates.

AcknowledgementsÐFinancial support for this work wasprovided by NSF (CMS 9503943). One of the authors(S.A.) gratefully acknowledges the University of Tehran,Tehran, Iran, for the initial support through a graduatescholarship. The MP35N alloy used in this study was pro-vided by SPS Technologies, Inc. (Jenkintown,Pennsylvania) and by Latrobe Steel Company (Pittsburgh,Pennsylvania). Helpful discussions with Steve Buzolits andMark L. Robinson are sincerely acknowledged.

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