influence of heat treatment time and temperature … of heat treatment time and temperature ... the...

CORROSION SCIENCE SECTION

055002-1 CORROSION—MAY 2011

Submitted for publication October 7, 2010; in revised form, December 29, 2010.

‡ Corresponding author. E-mail: [email protected]. * Department of Materials Science and Engineering, Lehigh Univer-

sity, 5 East Packer Ave., Bethlehem, PA 18015.

Influence of Heat Treatment Time and Temperature on the Microstructure and Corrosion Resistance of Cast Superaustenitic Stainless Steels

J.N. DuPont‡,* and J.D. Farren*

ABSTRACT

The influence of heat treatment time and temperature on the microstructure and corrosion resistance of cast stainless steel alloys CN3MN (UNS J94651) and CK3MCuN (UNS J93254) were investigated with a combination of microscopy tech-niques and immersion corrosion testing. The as-cast alloys exhibited significant residual microsegregation of Mo across the dendritic substructure and the presence of sigma (σ) phase in the interdendritic regions. Complete dissolution of the sigma phase and homogenization of Mo occurred after a heat treat-ment of 1,205°C for 4 h for castings with an average dendrite spacing of 35 µm. A heat treatment of only 1,150°C for 1 h was sufficient to produce similar results in welds with a smaller dendrite spacing of 10 µm. The use of the 1,205°C/4-h treatment restored the corrosion resistance of the cast alloys to a level that is comparable to the wrought counterpart alloys when evaluated using the ASTM G48 Method A test. This improvement is attributed to homogenization of Mo and disso-lution of the σ phase. Calculated curves are proposed to esti-mate the combination of time and temperature needed for effective heat treatment of castings with various dendrite spacings.

KEY WORDS: austenitic stainless steel, heat treating, homog-enization, localized corrosion, microsegregation

INTRODUCTION

Superaustenitic stainless steel castings are used in a variety of environments where good corrosion resis-

tance and toughness are required. Mo is a key alloy-ing element that is added for improved resistance to crevice and pitting corrosion. Recent work has sug-gested that Mo, in combination with N, promotes selective dissolution of iron at the surface and leads to Cr enrichment beneath the passive film, thus increasing corrosion resistance.1 A nitride layer has also been detected at the film-metal interface along with a ferrous molybdate (FeMoO4) layer in the outer regions of the passive film. These phases have been proposed to provide secondary kinetic barriers for further enhancement in corrosion performance. In wrought alloys, where the Mo is homogeneously dis-tributed, the protective surface film that develops provides excellent pitting and crevice corrosion resis-tance.2 However, the stability of the protective film is adversely affected in regions of Mo depletion that is commonly observed in castings and welds.3

Mo segregates to the liquid during solidification because of the relatively low solubility of Mo in aus-tenite. The inability of Mo to diffuse down the con-centration gradient as a result of the low diffusivity of Mo in austenite leaves the dendrite cores depleted in Mo.4 As a result, castings and welds are suscep-tible to preferential corrosive attack at the dendrite cores.3 The Mo enrichment in the liquid also leads to the formation of brittle interdendritic σ phase that has a deleterious influence on impact toughness.5 These deleterious effects potentially can be eliminated with a post-casting heat treatment designed to homoge-nize the Mo and dissolve the interdendritic σ phase. However, the influence of heat treatment time and temperature on the homogenization and dissolution

ISSN 0010-9312 (print), 1938-159X (online)11/000057/$5.00+$0.50/0 © 2011, NACE International


CORROSION—Vol. 67, No. 5 055002-2

behavior in these alloys has not yet been investigated. In addition, no reports are available that describe the influence of heat treatment on the resultant corro-sion behavior. The objective of this research was to determine the influence of heat treatment time and temperature on the microstructure and concomitant corrosion resistance of cast superaustenitic stainless steel alloys CN3MN (UNS J94651)(1) and CK3MCuN (UNS J93254). These alloys were selected because of their relatively high Mo concentrations and indus-trial experience that indicates significant differences in the corrosion performance of cast and counterpart wrought alloys. The results of this work are useful for designing industrial heat treatments that can be used to restore the corrosion resistance of cast alloys to a level similar to that of their wrought counterparts.

EXPERIMENTAL PROCEDURES

Alloy PreparationAlloys CN3MN and CK3MCuN were cast into sand

molds to produce samples that were approximately 25 by 25 by 3,050 mm in dimension. The compositions of the cast and counterpart wrought alloys are provided in Table 1. The numbers in parenthesis for the cast alloys are the allowable composition ranges accord-ing to ASTM A743.6 Alloys AL6XN† (UNS N08367) and 254SMO† (UNS S31254) are the counterpart wrought alloys for CN3MN and CK3MCuN, respectively. They were included in the experiment as a basis for com-parison. Samples for heat treating were extracted from the original castings and machined into 19- by 19- by 6.3-mm specimens. The samples were heat-treated at 1,150°C and 1,205°C for 1, 2, and 4 h in a laboratory tube furnace. Flowing Ar was utilized during the heat treatment to prevent surface oxidation. Samples from the wrought alloys were machined to the same dimen-sions from plate. Attempts were made initially to include heat treatment temperatures of 1,260°C and

1,315°C. However, subsequent examination of sam-ples exposed to these temperatures revealed a sec-ondary intergranular phase associated with liquation, indicating these temperatures were above the solidus temperatures of the alloys. Therefore, these tempera-tures were not evaluated further.

MICROSCOPY

Samples for microstructural characterization were mounted in thermosetting epoxy and polished to a 0.05-µm finish using standard metallographic tech-niques. Electrolytic etching was accomplished using a 10% sodium cyanide (NaCN) solution at 3 V to 5 V for 5 s to 10 s. Microstructural evaluation was performed on all samples using standard light optical microscopy (LOM) techniques. Secondary-phase fraction measure-ments were conducted using a Reichert-Jung† metal-lograph in combination with Leco† image analysis software. A minimum of 10 fields were used to measure the area fraction of the secondary phase, and the area fraction was assumed to equal the volume fraction.

Phase identification was conducted using electron back-scatter diffraction (EBSD) and x-ray energy-dis-persive spectroscopy (XEDS). The EBSD patterns and XEDS spectra were collected using an FEI† scanning electron microscope (SEM) equipped with both back-scattered electron and EDS detectors. HKL Flamenco† and HKL Spirit† software packages were used to col-lect and interpret the EBSD and EDS raw data. A 20-kV accelerating voltage was used for both tech-niques and a 60° tilt was required for the EBSD mea-surements. The EBSD patterns were processed using a band detection technique to calculate band spac-ings and angles and identify matches to known stan-dard patterns. The mean angular deviation (MAD) was used as an indication of agreement, with a MAD < 1° considered as confirmation for positive phase iden-tification.7 Electron probe microanalysis (EPMA) was performed using a JEOL 733† superprobe with inte-grated wavelength dispersive spectrometers (WDS). All measurements were made using a 20-kV accelerating voltage with a 35-nA beam current. The compositions

(1) UNS numbers are listed in Metals and Alloys in the Unified Num-bering System, published by the Society of Automotive Engineers (SAE International) and cosponsored by ASTM International.

† Trade name.

TABLE 1Chemical Compositions of Alloys(A)

Element CN3MN CK3MCuN AL6XN† 254 SMO†

C 0.02 (0.03 max.) 0.02 (0.025 max.) 0.02 0.02 Cr 21.45 (20 to 22) 19.91 (19.5 to 20.5) 21.44 20.02 Cu 0.06 (0.75 max) 0.55 (0.50 to 1.0) 0.19 0.56 Fe 45.92 (bal.) 52.39 (bal.) 47.22 53.23 Mn 0.47 (2 max.) 0.560 (1.2 max.) 0.25 0.55 Mo 6.40 (6 to 7) 6.30 (6 to 7) 6.30 6.40 N 0.23 (0.18 to 0.26) 0.224 (0.18 to 0.24) 0.23 0.215 Ni 24.4 (23.5 to 25.5) 18.94 (17.5 to 19.7) 23.80 18.3 P 0.013 (0.04 max.) 0.014 (0.045 max.) 0.028 0.029 S 0.007 (0.01 max.) 0.006 (0.01 max.) <0.001 <0.01 Si 0.99 (1 max.) 1.08 (1 max.) 0.30 0.67(A) All values in weight percent.



were measured using a ZAF data reduction technique that converts x-ray counts into weight percentages.8

CORROSION TESTING

After heat treating, samples for corrosion testing according to ASTM G48 Method A9 were prepared to an 80-grit surface finish. The samples were weighed to the nearest 0.001 g and immersed in an acidified ferric chloride solution (6% FeCl3 by mass, 1% hydro-chloride acid [HCl]) at 75°C for 72 h. After corrosion testing, the samples were scrubbed using a nylon brush, ultrasonically cleaned, and dried with etha-nol (C2H6O) to remove any debris from the pits. Total weight loss was used as an indicator of corrosion resistance.

RESULTS AND DISCUSSION

As-Cast MicrostructuresFigure 1 shows light optical photomicrographs of

each alloy in the as-cast condition. Also shown are the microstructures of autogenous welds placed on

each alloy (using the gas-tungsten arc process at 100 A, 9 V, and 3.0 mm/s travel speed). These welds were placed on the surface of each sample used for heat treating to assess the influence of dendrite spac-ing. The microstructure of each alloy consists of aus-tenite dendrites with an interdendritic eutectic-type constituent. Figure 2 shows a SEM image and EDS spectra of the interdendritic region. The spectra were acquired from the primary austenite (Figure 2[b]), eutectic-type austenite (Figure 2[d]), and secondary interdendritic phase (Figure 2[c]). Note that the inter-dendritic phase is enriched in Mo and Cr relative to the austenite. Typical EBSD patterns from each phase are shown in Figure 3. The calculated patterns are also shown and confirm that the matrix is austenite (Figures 3[b] and [d]) while the secondary phase is σ (Figure 3[c]). These results are consistent with the EDS spectra, since σ is known to be rich in Mo and Cr.4,10 The σ phase was the only secondary phase observed for all heat treatment conditions. Figure 4 shows an EPMA trace acquired across several den-drites in the as-cast CK3MCuN alloy. The Mo concen-tration varies from about 4.5 wt% in the dendrite

FIGURE 1. ([a] and [b]) Light optical photomicrographs of each alloy in the as-cast condition. Also shown are the ([c] and [d]) microstructures of autogenous welds placed on each alloy.



cores to about 9 wt% in the interdendritic regions. The Cr, Si, and Ni segregate in a similar direction (i.e., toward the liquid during solidification), while Fe seg-regates in the opposite direction. Results for alloy CN3MN showed essentially identical results.

Initial solidification of these alloys starts with pri-mary L → γ austenite, which leads to Mo rejection into the liquid.4,10 The process continues until the L → γ + σ eutectic-type composition is reached in the liquid, at which point solidification terminates with the forma-tion of an interdendritic γ/σ constituent. The residual microsegregation is expected in these alloys and is attributed to the low diffusivity of Mo in γ.4,10 This also leads to extensive enrichment of Mo in the liquid, which accounts for formation of the interdendritic σ

phase. It has been shown that the extent of microseg-regation and resultant secondary-phase formation is generally insensitive to the cooling rate for these alloys.11 Lower cooling rates during solidification pro-vides more time for back-diffusion during solidifica-tion, which, in itself, would reduce the extent of residual segregation. However, this effect is offset by larger dendrite spacings that are produced with slower cooling rates, so that the extent of microsegre-gation generally is not affected by cooling rate (within the range of rates typical for casting and arc welding).

Sigma Phase Dissolution KineticsFigure 5 shows light optical photomicrographs

of alloy CK3MCuN after heat treatment times of 1, 2,

FIGURE 2. (a) SEM image and EDS spectra of the interdendritic region from alloy CK3MCuN acquired from the (b) primary austenite, (c) secondary interdendritic phase, and (d) eutectic-type austenite.



and 4 h at 1,150°C. The sample heat-treated for 1 h also shows the weld. Note that dissolution is not com-plete within the casting at times up to 4 h at this tem-perature, but is complete in the weld after only 1 h. The dissolution observed in the weld indicates that 1,150°C is above the γ/σ solvus, but more time would be needed for dissolution in the casting because of

the larger dendrite spacing (average dendrite spac-ing of 35 µm in the casting compared to 10 µm in the weld). Figure 6 shows results for the casting for a temperature of 1,205°C. At this temperature, signifi-cant dissolution occurs after the 2-h heat treatment, and dissolution is complete after 4 h (the features vis-ible in Figure 6[c] are shrinkage porosity, not second-

FIGURE 3. Typical measured and calculated EBSD patterns from ([a], [b], and [d]) austenite and (c) σ phase within the interdendritic region.

FIGURE 4. EPMA trace acquired across several dendrites in as-cast CK3MCuN alloy showing (a) the full trace and (b) expanded view of Mo and Si concentrations.

(a) (b)



ary phase). Identical results were observed for alloy CN3MN.

Figure 7 shows the variation in σ phase content as a function of heat-treating time and temperature for each alloy. Also shown in each plot is the com-puted dissolution behavior of σ phase as determined by the model proposed by Singh and Flemings.12

V V

C CC

DttV VtV V MC CMC C= +V V= +V V

V V

V VV V= +V V

V V= +V V

V V= +V V

V V= +V VV V= +V VV V= +V VV V

V VV V

V VV V= +V V

V V= +V VV V= +V V

V V= +V V

V V= +V V

V V= +V VV V= +V V

V V= +V V

V VV VV V= +V VV V= +V VV VV V

V VV VV V= +V VV V= +V V

V V= +V VV V= +V V

0= +0= +V V= +V V0V V= +V V 0

2

02

C C–C Cexp –

p –

p –

p –

p –

p –σ

π4 l

–

C C–C C–C

M 0C CM 0C C–C C–M 0–C C–

σ (1)

Singh and Flemings derived this expression for a binary alloy containing microsegregation and non-equilibrium eutectic. They used this expression to investigate dissolution kinetics in multicomponent Al alloys by treating the alloys as an Al-Cu binary sys-tem. This expression was adopted to the current work by treating the alloys as a γ-Mo binary system, since Mo is the primary element that segregates during solidification and leads to formation of the Mo-rich σ phase while Ni, Cr, and Si do not segregate strongly (as shown in Figure 4). In this case, Vt is the volume fraction of σ after dissolution time t, V0 is the initial volume fraction of σ, CM is the maximum solid solubil-

ity of Mo in γ, C0 is the nominal Mo concentration, Cσ is the Mo concentration of the σ phase, l0 is half the dendrite spacing, and D is the diffusivity of Mo in γ. V0 and l0 were measured for each alloy using image analysis techniques. The value for Cσ was calculated at 9 wt% Mo with Thermo-Calc† software using the CALPHAD method with the Fe database.13 This value is in good agreement with the EPMA measurements that show the maximum concentration of Mo in the primary austenite was also ~9 wt% Mo. The value for Cσ was determined in a previous study as 25.9 wt% Mo using EPMA.4 Values of DMo were calculated from the diffusivity data provided by Malik, et al.,14 in which Do = 1.1 × 10–4 m2s–1 and Q = 275 kJ mol–1.

There is reasonable agreement between the calcu-lated and measured σ phase contents in Figure 7 con-sidering the difficulty in measuring secondary phase fractions at these low values and that a binary dis-solution model is applied to these multi-component alloys. The measurements and calculations each indi-cate that dissolution will not be complete at 1,150°C up to 4 h, while a 4-h to 5-h exposure at 1,205°C dis-solves the σ phase. The calculations also indicate that complete dissolution should occur in the weld at 1,150°C for 1 h due to the reduced dendrite spacing of

FIGURE 5. Light optical photomicrographs of alloy CK3MCuN after heat treatment times of (a) 1, (b) 2, and (c) 4 h at 1,150°C. The sample heat-treated for 1 h also shows the (d) weld.



10 µm, which is also consistent with the observations (Figure 5).

Homogenization KineticsFigure 8 shows typical EPMA Mo concentra-

tion profiles acquired on alloy CK3MCuN after heat treatment conditions of 1,150°C/4 h (Figure 8[a])

and 1,205°C/4 h (Figure 8[b]). Significant residual microsegregation persists after the 1,150°C/4-h treat-ment, whereas the residual segregation is essen-tially eliminated after the 1,205°C/4-h treatment. The EPMA data was used to determine the index of resid-ual segregation of Mo for each heat treatment condi-tion via:

δ

θ θ

iM m

M m

C Cθ θC Cθ θM mC CM m

C CM mC CM m

= C C–C CC C–C C0 0C C0 0C C

(2)

where CθM and Cθ

m represent the maximum and mini-mum Mo concentrations after a given homogeniza-tion time, and C0

M and C0m represent the maximum

and minimum initial Mo concentrations, respectively. Complete homogeneity was confirmed using a homo-geneity criteria, which determines whether the raw, uncorrected x-ray intensities collected for each ele-ment fall within the statistical limits of a homogenous sample. The homogeneity criteria is given as:7

N NN NN N N N Ni– 3N N– 3N N 3≤ ≤N N≤ ≤N NN NiN N≤ ≤N NiN N + (3)

where Ni is the measured Mo x-ray intensity of a given measurement and N

– is the average Mo x-ray intensity

of all measurements. Heat treatment conditions in which all x-ray intensity measurements fell within the sample homogeneity criteria are statistically homoge-nous and assigned a d value of zero.

The variation in d for Mo with time and tempera-ture is shown for each alloy in Figure 9 (EPMA mea-surements were not conducted on all heat treatment conditions). Note that ~40% to 70% of the residual segregation of Mo still persists after the 1,150°C/4-h treatment for the casting. By contrast, the weld is nearly completely homogenized (d ~ 0.1) after a 1,150°C/1-h treatment as a result of the finer den-drite spacing. Complete homogenization of the cast-ing was only observed for the 1,205°C/4-h treatment. Also shown in Figure 9 is the calculated variation in d with time and temperature according to the sim-ple homogenization model originally proposed by Kattamis and Flemings15 for a single-phase alloy:

δ π

iDtl

=

exp –

p –p –

p –

2

0 (4)

where the values are as previously defined. The measured homogenization kinetics are significantly slower than those calculated using Equation (4). This is to be expected since Equation (4) assumes a single-phase alloy and therefore does not account for the influence of sigma phase on the required homogeniza-tion time. The presence of sigma phase is significant because it provides a source of Mo diffusion into the primary austenite as it dissolves, and this is not accounted for in the single-phase homogenization

FIGURE 6. Light optical photomicrographs of alloy CK3MCuN after heat treatment times of (a) 1, (b) 2, and (c) 4 h at 1,205°C.



(a)

(a)

(a)

(b)

(b)

(b)

FIGURE 7. Variation in sigma phase content as a function of heat-treating time and temperature for (a) CN3MN and (b) CK3MCuN.

FIGURE 8. Typical EPMA composition traces acquired on alloy CK3MCuN showing the distribution of Mo after heat treatment at (a) 1,150°C for 4 h and (b) 1,205°C for 4 h.

FIGURE 9. Variation in index of residual segregation (d) with time and temperature for each alloy.



model above. Therefore, complete homogenization is not attainable until the sigma phase is completely dissolved.

Corrosion ResistanceAs previously described, the relatively poor corro-

sion resistance of castings and welds in these alloys has been attributed to Mo microsegregation. With this in mind, Figure 10 shows the variation in weight loss from the ASTM G48 test as a function of the index of residual segregation for the cast samples (welds were not corrosion tested in this program). Corrosion occurred by a combination of general cor-rosion and pitting. The measured weight loss for each counterpart wrought alloy is also shown for compari-son. There is a clear correlation between the corro-

sion performance and index of residual segregation. In particular, the corrosion resistance in this test is comparable to that of the wrought alloys for a heat treatment of 1,205°C/4 h. Figure 11 shows corro-sion results for all the heat treatment conditions. (EPMA measurements were not conducted on all the heat-treated samples, so d values are not provided for all the heat treatment conditions in Figure 10.) The characterization results demonstrated that the 1,205°C/4-h treatment was effective at eliminating residual microsegregation and dissolving the second-ary sigma phase. The improvement in corrosion resis-tance is attributed to these favorable microstructural changes.

These results are significant with regard to typi-cally used industrial heat treatments. For example,

(a)

(a)

(b)

(b)

FIGURE 10. Mass loss vs. index of residual segregation for (a) CN3MN and (b) CK3MCuN.

FIGURE 11. Variation in weight loss from the ASTM G48 test as a function of sample condition for (a) CN3MN and (b) CK3MCuN.



the pertinent specification for these alloys6 requires a minimum postcasting heat treatment of 1,150°C/1 h. However, the results shown here clearly demonstrate that, for the typical casting conditions and dendrite spacings evaluated in this work, this heat treatment will not promote significant homogenization or disso-lution. Higher temperatures and longer times would be needed for significant improvement.

The required combination of time and temper-ature for complete dissolution will be affected by the cooling rate and resultant dendrite spacing. As described above, dissolution of the interdendritic σ phase requires longer times than homogenization of Mo within the dendrites. In addition, reasonable agreement was observed between the volume frac-tion of the σ phase measured experimentally and cal-culated with Equation (1). Therefore, the dissolution model in Equation (1) can be used to estimate the times and temperatures required for effective heat treatment as a function of cooling rate and resultant dendrite spacing. The variation in dendrite spacing (λ) with cooling rate (ε) for these alloys can be esti-mated by:16

λ ελ ε=λ ελ ε8λ ε 0 33λ ε0λ ε– .0 3– .0 3 (5)

where λ is measured in µm and ε is calculated in °C/s. Once λ is known by means of the cooling rate, the time and temperature required for dissolution can be estimated using Equation (1). Figure 12 illustrates the results of such calculations for each alloy. It is recognized that these results are only approximations that are subject to the accuracy and limitations of the dissolution model. Nevertheless, these curves are use-ful because they permit estimation of the required times and temperatures that are needed for dissolu-tion as a function of the casting cooling rate, which is commonly determined using commercially available heat flow software.

CONCLUSIONS

The influence of heat treatment time and temper-ature on the microstructure and corrosion resistance of cast stainless steel alloys CN3MN and CK3MCuN were investigated with a combination of microscopy techniques and immersion corrosion testing. The fol-lowing conclusions can be drawn from this work.❖ In the as-cast condition, the alloys exhibit primary austenite dendrites with residual microsegregation of Mo and interdendritic σ phase. The formation of σ phase is attributed to liquid enrichment in Mo during solidification.❖ Complete homogenization of Mo and dissolution of σ phase can be achieved in castings with an average dendrite spacing of 35 µm with a 1,205°C/4-h heat treatment.❖ The 1,205°C/4-h heat treatment restored the corro-sion resistance of each alloy to a level that is compa-rable to the wrought counterpart materials.❖ Heat treatment curves are proposed that provide estimates of the times and temperatures required to restore the corrosion resistance of cast alloys with various dendrite spacings.

ACKNOWLEDGMENTS

This article is based on work supported by the U.S. Department of Energy under Award no. DF-FC36-04GO14230. Any findings, opinions, and conclusions or recommendations expressed in this article are those of the authors and do not necessarily reflect the views of the Department of Energy. The authors grate-fully acknowledge members of the Steel Founders’ Society of America (SFSA) High Alloy Research Com-mittee for supplying materials and helpful technical guidance. Special thanks go to R. Bird (Stainless Foundry) and M. Blair (SFSA) for insightful technical discussions during the course of this research.

(a) (b)FIGURE 12. Required heat treatment times and temperatures for various dendrite arm spacings for (a) CN3MN and (b) CK3MCuN.



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influence of heat treatment time and temperature … of heat treatment time and temperature ... the...

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