The cyclic stress-strain behavior of nickel-base superalloys—II. Single crystals

~cta metall. Vol. 36, No. 2, PP. 283-290, 1988 Printed in Great Britain. All rights reserved

o@Ol-6160/88 $3.00+ 0.00 Copyright 0 1988Pergamon Journals Ltd

THE CYCLIC STRESS-STRAIN NICKEL-BASE SUPERALLOYS-II.

BEHAVIOR OF SINGLE CRYSTALS

L. G. FRITZEMEIERT and J. K. TIEN Center for Strategic Materials, Henry Krumb School of Mines, Columbia University, New York, NY 10027, U.S.A. (Received 19 May 1984; in revised form 18 May 1987)

Abstract-The cyclic stress-strain behavior of the single crystal nickel-base superalloy Mar-M200 has been investigated at ambient temperature and 843°C under constant strain rate cycling at 10e2 and 10-‘/s. At ambient temperature, the cyclic stress-strain curve exhibits a clear strain rate independent plateau regime similar to that observed for f.c.c. metal single crystals and other coherent precipitate strengthened systems. At elevated temperature, the cyclic stress-strain curves are reduced to power law functions with the higher strain rate curve exhibiting a higher slope, or work hardening exponent than the lower strain rate curve. The forms of these curves, the cyclic hardening behavior and details of hysteresis behavior are rationalized with respect to substructural observations. At ambient temperature, the substructure consists of widely spaced, well defined slip bands which shear the y’ precipitates and extend completely through the crystal on { Ill} planes. At 843°C and 10m2/s,the slip bands initially shear the y’ which then reform around the slip bands. In addition, a more homogeneous tangled structure which cages the y’ and provides a work hardening component is found. The lower strain rate at elevated temperatures produces a homogeneous substructure throughout the crystal with evidence of less well defined bands. These results are compared and contrasted with the results of a companion study (Part I) on a representative polycrystalline superalloy. Rt%um~Nous avons Ctudit le comportement cyclique contraintedeformation du superalliage monocristallin a base nickel Mar-M200, I la temperature ambiante et a 843C, pour une vitesse de deformation constante et des frequencies de 10-‘/s et 10-‘/s. A la temperature ambiante, la courbe cyclique contrainte-diformation prtsente un regime de palier net independant de la vitesse de deformation, semblable a celui que l’on observe pour les monocristaux de mirtaux cfc et d’autres systemes durcis par des precipites coherents. A haute temperature, les courbes cycliques contrainte-&formation se rambnent a des fonctions en loi de puissance, la courbe correspondant a la vitesse de deformation la plus &levee presentant une pente (ou un exposant de durcissement) plus prononcee que la courbe correspondant a la vitesse de deformation la plus basse. Nous avons rationalis la forme de ces courbes, le durcissement cyclique et les details de l’hystertsis en tenant compte des sous-structures observtes. A la temperature ambiante, la sous-structure est constituee par des bandes de glissement bien definies et largement espacees, qui cisaillent les pricipites y ’ et traversent entierement la cristal sur des plans {111). A 843C, et pour IO-‘/s, les bandes de glissement commencent par cisailler les precipitis y’ qui se reforment ensuite autour d’elles. En outre, nous avons trouvt une structure enchevdtrbe plus homogene qui emprisonne les precipitis y’ et provoque une composante de durcissement. Aux temperatures Clevtes, la vitesse de deformation la plus basse produit une sous-structure homogene a l’tchelle du cristal, avev des bandes de glissement moins bien definies. Nous comparons ces rtsultats, et nous les confrontons avec ceux dune etude similaire. Zusammenfassung-Das

zyklische Spannungs-Dehnungsverhalten eines Einkristalles der Superlegierung Mar-M200 auf Nickelbasis wurde bei Raumtemperatur und bei 843°C unter Zyklen bei konstanter Dehnungsrate von 1O-2/s und 10-‘/s untersucht. Bei Raumtemperatur weist die zyklische Spannungs-Dehnungskurve ein deutliches, von der Dehnungsrate unabhingiges Plateau auf, der dem bei kfz. Metallkristallen und in anderen ausscheidungsgehirteten Systemen beobachteten ahnelt. Bei erhiihten Temperaturen sind die Spannungs-Dehnungskurven auf Potenzfunktionen reduziert, wobei die Kurve fiir die gr68ere Dehnungsrate eine grol3ere Steigung, oder einen griil3eren Verfestigungskoeflizienten, als die fur niedrige Dehnungsrate aufweist. Die Form dieser Kurven, das zyklische Verfestigungsverhalten und Einzelheiten des Hystereseverhaltens werden im Hinblick auf die Versetzungsstruktur behandelt. Bei Raumtemperatur besteht die Versetzungsstruktur aus gut definierten Gleitblndem mit groBem gegenseitigem Abstand, die die y’-Ausscheidungen scheren und sich vollstiindig durch den K&all auf { 11I}-Ebenen ausdehnen. Bei 843°C und 10e2/s scheren die Gleitblnder die y’-Ausscheidungen anfangs; diese bilden sich jedoch in der NThe der Gleitblnder wieder zuriick. Dariiberhinaus wird eine homogenere verkniiuelte Struktur beobachtet, die die y ‘-Ausscheidungen umgibt und damit zur Verfestigung beitrlgt. Bei der niedrigeren Dehnungsrate bei hoherer Temperatur entsteht eine homogene Substruktur im gesamten Kristall, Anzeichen von weniger gut definierten Gleitbiindern finden sich. Diese Ergebnisse werden mit denen einer begleitenden Untersuchung (Teil I) an einer reprlsentativen polykristallinen Superlegierung verglichen. 1. INTRODUCTION It was found in a companion study (hereinafter

called

tPresent address: Rocketdyne Division, Rockwell International, 6633 Canoga Avenue, Canoga Park, CA 91303, U.S.A.

Part I) to this study that the behavior of a typical polycrystalline onic 115, at ambient temperature behavior exhibited by f.c.c. metals elevated temperature, the cyclic 283

cyclic stress-strain superalloy, Nimwas similar to the such as copper. At stress-strain curve

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Table I. Mar-M200 single crystal chemistry in wt% Ni-Bal. w-12.0

cr-9.0 CL-l.0

vo-IO c- < 0.01

Ti-2.0

Al-5.0

Mar-M200 sin& crvstal heat treatment 4 h at 1250°C 4h at 1080°C 32h at 871°C

reduces to an approximately power law function with no apparent plateau behavior, although cyclic hardening behavior and details of hysteresis loop shape were relatively unchanged with temperature. The strong plateau behavior generally associated with persistent slip band formation, such as developed in this alloy at ambient temperature, was suppressed by a more homogeneous dislocation substructure forced to occur due to grain boundary strain accommodation effects. This paper (Part II) contains the results of a study of the cyclic stress-strain behavior of a representative, nickel-base superalloy single crystal, Mar-M200, allowing the avoidance of the grain boundary effects in such studies. In addition, the effect of varied strain rate on cyclic stress-strain response has been included for study.

2. EXPERIMENTAL PROCEDURE AND STRATEGY 2.1. Materials and specimen preparation Ideally this study should be conducted using a single alloy composition comparable to the polycrystalline chemistry from Part I. In practice the requirements for single crystal chemistries are different from the needs of a useful polyscrystalline nickel-base superalloy system. Since no servicable system is presently available in both forms a complete alloy development program would have been required to fulfill this need. An available representative single crystal system was then chosen for this study. A Mar-M200 single crystal casting 76 mm in diameter and 150 mm in length was obtained. Mar-M200 is a

---J

n

solutionizing step/air cool simulated coating step/air cool urecioitate aging/air cool

900

MAR-

M200

Strain

amplitude

2

0

20

smgle

crysfal

5.4x

843

C

IO“

40

60

60

100

Cycles Fig. 2. Typical Mar-M200 cyclic hardening curve representative of the CSSC plateau. The total strain amplitude is 5.4 x 10-2

tested, carbide free, commercially available single crystal alloy. Chemistry and heat treatment schedules are presented in Table 1. An example of the microstructure is presented in Fig. 1. Approximately 60 vol.% of the alloy is y’ with an average edge length of about 0.25 pm. All single crystal specimens were taken from blanks cut parallel to the cube axis growth direction of the single crystal. Subsequent to heat treatment, specimen blanks were oriented on the (001) axis by back-Laue X-ray diffraction with a 0.5 deg resolution. Specimens subsequently machined from these blanks were then within 2 deg of the cube axis. The specimens initially employed for the testing were a button-headed design as described in Part I. It was discovered shortly after beginning testing of the single crystal specimens that premature failure occurred at the extensometer ridges. the bulk of the testing was then conducted on smooth bar samples with extensometry attached near the inside faces of the specimen grips. Testing procedures are described in detail in Part I. Testing was expanded in this portion of the study to include two strain rates, lo-* and 10-‘/s. At the elevated temperature, the higher of these strain rates has been shown to be near the region where mechanical properties become strain rate independent while the lower strain rate is within the strain rate dependent regime in monotonic deformation [I, 21. 3.

EXPERIMENTAL RESULTS

3.1. Mechanical data

Fig.

I. Mar-M200

microstructure.

A typical cyclic hardening curve for the Mar-M200 single crystal tests is given in Fig. 2. Again, as discussed in Part I, there is no difference in cyclic hardening behaviour between ambient temperature and elevated temperature. Initially a slight degree of hardening occurs within the first ten to 50cycles, followed by saturation.

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Table 2. Elevated temperature power law fit parameters Au/2 = a,.(Ac,/2)“’ ok. (MPa)

n’

Mar-MZOO, 10-*/s 7800 Mar-MZOO, 10-‘/s 2412

0.327 0.168

2

OF Ni-BASE SUPERALLOYS-

1000 r

MAR-M200

285

smgle crysto,

/ _Y

900

z -

800

9 e

700

6 $

600

4 E

500

+ Ambient temperoWe

The ambient temperature Mar-M200 cyclic stress-strain curve is presented in Fig. 3. This curve clearly exhibits a strain amplitude independent plateau behavior with high work hardening taking place below strain amplitudes of approximately 5 x lo-‘, similar to a previously reported value for a coarse grained superalloy [3], and a plateau stress of approximately 850 MPa. There is no third stage evident in this curve. Cyclic stress-strain curves for Mar-M200 at 843°C and 10m2/s and 10m3/s strain rates are also given in Fig. 3. It is immediately evident that the IO-‘/s curve exhibits a higher slope than the 10W3/s curve. This higher cyclic work hardening exponent causes the curves to cross over at intermediate strain amplitudes. These curves lie well below the ambient temperature curve for low and intermediate strain amplitudes. The finite slopes of the elevated temperature curves cause them to approach the ambient temperature plateau at higher strain amplitudes. Power law equation fits to the elevated temperature curves are given in Table 2. Comparison of the cyclic and monotonic 0.2% yields is given in Table 3. At ambient temperature the cyclic yield is significantly higher than the monotonic value. The elevated temperature monotonic and cyclic yields are in closer agreement, with a lower strain rate producing a lower yield than a high strain rate. Resisting stress (a,), or hysteretic yield, as defined in Part I, curves for Mar-M200 are presented in Fig. 4. The curves are polynomial best fits to the data points. At ambient temperature the curve is nearly flat with a slight increase in or as strain amplitude increases. Both elevated temperature curves begin at low values, increasing to an apparent saturation level with strain amplitude. The elevated temperature curves both lie below the ambient temperature curve with the higher strain rate curve the higher of the two.

?I

. a43 oc

0843’C

400 300

II

10-U

10-s

Plastic

Fig. 3. Mar-M200

Fracture behavior was found to be unchanged over the range of the plateau region and surfaces characteristic of the failures of the Mar-M200 single crystals within this region are presented in Fig. 5. At ambient temperature and at 10e2/s strain rate at the elevated temperature, the fracture is highly crystallographic along the { 111) planes approximately 54 deg to the

900

-

_

800

r

i+i

500

A

843°C

10--2/s IO~~‘/S

0

0.2% yield stresses

Cyclic-1000 900 MPa 850 MPa

MPa

Monotome 954 MPa 780 MPa

cyclic stress-strain

MAR- M200

A

curves

res,sf,ng S,WSS

0

24°C

a

.943oc; 943oc

A

I/I

temperature

10-z

amplitude

(001) stress axis. Specific initiation sites are difficult to pinpoint due to obliteration of fine details by rubbing during the strain cycles. Fractures at 843°C and lO-3/s exhibit a combination of Stage I facets in a macroscopically Stage II more ductile surface. Representative transmission micrographs of the substructures developed during the ambient temperature fatigue of Mar-M200 single crystals are presented in Fig. 6. The substructure produced by cycling above the lower plateau limit is extremely heterogeneous and consists of sharp, well defined slip bands along the { 11 l} planes. The material away from these slip bands exhibits no indication of deformation. The slip bands have shear,ed the y’ precipitates encountered as they traverse the diameter of the specimen. It is noteworthy that substructure was encountered in only about one of every ten foils prepared for microscopy. Substructures observed in the material deformed at lO-2/s at 843°C are presented in Fig. 7. Again numerous foils had to be prepared to find substructural evidence. The substructure varies from slip bands which appear to follow a somewhat tortuous path roughly parallel to the { 111) slip

tj

Ambient

10-s

strain

0

3.2. Fractography and transmission microscopy

Table 3. Mar-M200

10-s/ I 10-2/s

0.0005

0.001 0.0015 0.002 Plastic

MPa

Fig. 4. Mar-M200

1o-2/s ; 10-3/s

strain

amplitude

resisting stress as a function strain amplitude.

of plastic

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Fig. 5. Characteristic fracture surfaces from the intermediate strain range region for Mar-M200 crystals representing (a) ambient temperature and IO-*/s at 843°C and (b) 10m3/sat 843°C

Fig, 6. Representative

TEM microscopy for Mar-M2~ fatigued at ambient temperature within the intermediate strain range regime.

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Fig. 7. Representative TEM microscopy for Mar-M2~ fatigued at IO-‘/s and 843°C in the i~t~~mediate strain range regime.

traces but between y’ precipitates, to a looser more homogeneous structure caging the y’, At 843°C and 1W3/s the structure is typically as found in Fig. 8. All foils prepared exhibited the loose homogeneous tangles caging the y ’ precipitates. However, some evidence of planar slip and slip bands similar to those found at 1O-2/s and 843°C was found. 4. DISCUSSiON OF RESULTS It is useful to briefly summa~ze and contrast the important trends observed in the results from both Part I and Part II, representing polycrystals and single crystals, respectively. 1. The cyclic hardening curves exhibit slight hardening to saturation. 2. The ambient temperature cyclic stressstrain curve of the single crystal alloy exhibits clear plateau behavior, in contrast to the less distinct ~lycrystalline behaviour reported in Part I.

3. The elevated temperature cyclic stressstrain curves of both superalloys are of power law form. The higher strain rate curve exhibits a higher slope than the low strain rate curve for the single crystal case, 4. The single crystal ambient temperature resisting stress curve is relatively insensitive to strain amplitude. 5. The single crystal elevated temperature resisting stress values are low at low strain amplitudes, increasing as strain amplitude increases, as opposed to the decreasing values exhibited by the polycrystahine curves. 6. Cyclic yield values are similar to values from monotonic tests for both single crystal and polycrystal. The obvious plateau behavior in the cyclic stress-strain curve strongly suggests that PSB’s are the dominant deformation mode in Mar-M200 single crystals at ambient temperature. The resisting stress analysis indicates that these slip bands must either not shear the 1” phase or that another hardening

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Fig. 8. Representative TEM microscopy for Mar-M200 fatigued at 10B3/s and 843°C in the intermediate strain range regime. mechanism must take place such that these PSB’s do not substantially soften the material. The slight amount of hardening observed in the cyclic hardening curves indicates that the latter suggestion may hold true, A previous study[4] has indicated that stress cycling in nickel-base superalloys may cause cross slip even at ambient temperatures. This mechanism will cause a slight amount of work hardening within the slip bands whether or not y’ precipitates are sheared, thus accounting for the slight cyclic hardening behaviour and apparent retention of deformation resistance, consistent with cyclic stress-strain curve behavior. Slip bands in the Mar-M200 monocrystals at ambient temperature were found to be very well defined and widely spaced, as observed in the TEM foils, and extended completely through the crystal, as indicated by slip lines on the specimen surface. The slight hardening observed in the superalloy may be further explained as being due to the intersection of slip bands from competing (111) planes. It is possible that this hardening would not be observed for crystals in single slip orientation. The slip bands are

observed to shear the y’ on a fine scale with no indication that the process is exacerbated at increased strain amplitudes. The relatively long mean free path within the PSB’s allows dislocation motion with no requirement of cross slip and therefore no further y’ disruption should occur once the slip bands are established. The elevated temperature mechanical response suggests that at both strain rates the substructure is more homogeneous than at ambient temperature. The absence of a plateau and finite slopes in the cyclic stress-strain curve indicate that slip bands are not the primary substructure under these conditions. The increase of 6, (resisting stress) with strain amplitude supports the presence of a more homogeneous, tangled substructure at elevated temperature. The observed substructures at elevated temperature are somewhat different from the ambient temperature case, consistent with the mechanical response analysis. At IO-*/s which is near the strain rate independent portion of the 843°C monotonic yield stress vs strain rate curve, the substructure consists of slip bands and areas of homogeneous

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tangled substructure. The homogeneous portion of the substructure occurs in the matrix channels between the y ’ precipitates and can partially account for the increase in or with strain amplitude. It it also significant that the slip bands no longer appear purely planar under these conditions and appear internally less dense than those observed at ambient temperature. Diffusive reconstruction of the previously sheared y’ appears to have taken place, causing slip bands to be contained in the matrix and avoiding further microstructural disordering in the precipitate phase. At IO-‘/s and 843°C the substructure is much more with loosely tangled substructure homogenous, throughout each specimen and an apparently lower PSB density than the higher strain rate samples. This is again consistent with all observed mechanical behavior. As strain amplitude increases, dislocation density increases, resisting stress increases and total stress amplitude follows suit. The increased cross slip allowed under these conditions can also contribute to an increase in work hardening relative to the higher strain rate, providing explanation for the higher saturation stresses at low strain amplitudes. It is interesting that the structures developed in the polycrystalline superalloy discussed in Part I are significantly different from those developed in the monocrystals. However not all of this difference can be attributed to the presence, or absence, of grain boundaries. The necessity of using two different superalloy chemistries introduces additional variables into the experiment. Some details must undoubtedly be influenced by the differences in solid solution, or matrix strengthening, element content, y’ volume fraction and morphology and microstructural stability at elevated temperature. It is unfortunate that these problems cannot be overcome at this time, but it is felt that these effects are minor compared to the general trends discussed. The stress for the onset of the low slope region in the ambient temperature polycrystalline superalloy cyclic stress-strain curve and the plateau stress in the monocrystal ambient temperature curve are of interest since these values coincided with the fail-safe fatigue limit in the previous cyclic stress-strain studies. In both cases the equivalent stress amplitude is approximately 0.7 to 0.8 of the monotonic yield point, much higher than-the generally reported fatigue limits of near 0.5 of the yield for this alloy class[S]. The break points reported here are dependent only on substructural development whereas fatigue limits achieved by normal means are, by definition, failure limited. The high strengths and low ductilities of y’ strengthened systems make them highly susceptible to crack initiation from inclusions and other inhomogeneities. As illustrated by our experience with failure initiated at the extensometer ridges in the monocrystal specimens, anisotropy and heterogeneous slip cause the single crystals to be especially sensitive to stress concentration sites. In

OF Ni-BASE

SUPERALLOYS-

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this sense, the break points in the cyclic stress-strain curves of superalloys, even if found, can then be useful only as qualitative indicators of fail-safe limits in fatigue. 5. CONCLUDING

REMARKS

The cyclic stress-strain behavior of the single crystal nickel-base superalloy Mar-M200 has been investigated and compared with the results from a companion (Part I) polycrystalline superalloy study. The ambient temperature cyclic stress-strain curves produced here are similar in form to the respective f.c.c. metal (copper) counterparts. In particular, the single crystal curve exhibits clear plateau behavior while the polycrystalline curve indicates less clearly differentiated three stage behavior. At elevated temperature, the cyclic stress-strain curves are reduced to power law functions. Both materials exhibit cyclic hardening behavior with the polycrystalline material hardening to a greater extent. Significant differences exist in the details of stress-strain hysteresis loops from the two superalloys. The hysteretic yield of the polycrystalline material decreases with strain amplitude at both ambient and elevated temperatures. The ambient temperature hysteretic yield of Mar-M200 is roughly constant, while at elevated temperature it increases with strain amplitude. These findings are confirmed through, and are consistent with, TEM substructural observation. Deformation in the polycrystalline superalloy occurs in slip bands in the specimen interior and in dense tangles caused by grain boundary strain accommodation. Thus, although the slip bands disrupt the strengthening y’ precipitates and reduce the hysteretic yield, the peak cyclic stress remains high due to pile ups in the slip bands at grain boundaries and in the tangled structure. At ambient temperature the single crystal superalloy deforms exclusively through widely spaced, sharp slip bands which extend completely through the crystal on { 111} planes. Once these slip bands are set up no significant changes in material integrity occur, therefore accounting for the cyclic stress-strain curve plateau and stable hysteretic yield. At elevated temperature, the slip bands become less defined as the sheared y’ precipitates can reform around them. In addition, a more homogeneous, y’-caging substructure evolves which provides a work hardening component that suppresses plateau behavior in the cyclic stress-strain curve and accounts for the increasing hysteretic yield. A lower strain rate at this temperature promotes formation of the homogeneous substructure at the expense of slip bands.

Acknowledgements-The authors wish to thank the Exxon Research and Engineering center for partial support of this work, Campbell Laird for his helpful discussion and Maury Gel1 for providing the single crystal material.

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REFERENCES

OF Ni-BASE SUPERALLOYS-11

3. J. N. Vincent and L. Remy, in Fracture and the Role of Microslructu~ re, Proc. Fourrh ECF Conf. (1982).

1. G. R. Leverant, M. Gel1 and S. W. Hopkins, Mater. Sri. Engng 8, 125 (1971). 2. R. R. Jensen and J. K. Tien, in Metallugical Treatises

(edited by J. K. Tien and J. F. Elliot), pp. 529-550. T.M.S.-A.I.M.E. (1981).

4. G. L. Chen, L. G. Fritzemeier, X. Xie and J. K. Tien, ,?A lo<, 11o*‘), Metall. Trany.P xdn, .,_j. L.,VL,. 5. J. K. Tien and B. H. Kear, Scripta metall. 6, 303 (1972).