The Effects of Temperature and Hold Time on Dynamic Strain Aging in a

Nickel Based Superalloy

C. Cornet*,a,C. Lupton*, C. Stöcker**, K. Wackermann**, H.-J. Christ**M. Hardy*** and J.

Tong*, b

*Mechanical Behaviour of Materials Laboratory, School of Engineering, University of

Portsmouth, Portsmouth, PO1 3DJ, UK

**Institut für Werkstofftechnik, Department Maschinenbau, Universität Siegen,

Paul-Bonatz-Str. 9-11, 57076 Siegen, Germany

*** Rolls-Royce plc, Derby, DE24 8BJ, UK

Abstract

Serrations known as Portevin Le-Chatelier have been observed in a nickel-based superalloy

RR1000, which drastically affected the stress relaxation behaviour of the material. Further

experiments have been carried out over a range of temperatures and the mechanism of dynamic

strain aging has been studied using TEM. In addition to serrations, the alloy also displayed strain-

rate insensitivity and increased hardening with a temperature range from 300°C to 750°C.

Keywords:

Dynamic Strain Aging, High Temperature Fatigue, Dislocation Arrangements, Dwell Time,

Nickel Alloy

aNow at Airbus,

bCorresponding author. Tel.: +44 23 92842326; fax: +44 23 92842351. E-mail address: jie.tong@port.ac.uk (J. Tong).

1. Introduction

Dynamic strain aging (DSA) is one of the hardening mechanisms affecting many materials, which

manifests itself by “jerky” or serrated plastic flows and inhomogeneous yielding. These

instabilities are referred to as Portevin Le-Chatelier effect [1], corresponding to temperature

dependent strain localisation within a specific range of strain and strain rate as a consequence of

DSA [2]. DSA may be attributed to the interaction between diffusing point defects, such as solute

atoms and mobile dislocations, during plastic flow [3-5]. Serrations may be attributed to a

repeated snowballing effect of locking and unlocking of dislocations by these defects, leading to

heterogeneous deformation with an alternating increase/decrease in plastic activities, despite of

a monotonic straining.

DSA has been identified in several Ni-based superalloys including Inconel 718 [6], Udimet

720 [7], Udimet 720Li [8] and Nimonic [9], as well as in aluminium and steel alloys. In most of

these alloys, DSA is understood to be the result of carbon atoms diffusing along the dislocation

cores and stopping them from progressing further. In Waspaloy, Hayes and Hayes [10, 11]

suggested that the DSA mechanism resulted from the interaction between the carbon

atmosphere and the strengthening precipitates.

In this paper, we report a study of the DSA behaviour in a coarse grain variant of Alloy

RR1000 (CGRR1000), which was developed via power metallurgy route for turbine discs in the

latest aero-engines. Turbine discs are fracture critical components such that failure of such a

component can lead to the loss of the aircraft. Turbine discs are typically subjected to variable

temperatures from up to 750°C in the periphery to about 650°C at the centre. It is important to

understand the effects of temperature on the mechanical behaviour of the alloy so that accurate

life prediction can be made. To this end, experiments have been carried out under multiple

hardening and relaxation (MHR), stepwise constant strain rate in tension and strain rate-

controlled cyclic loading conditions, in order to capture most of the critical service loading

conditions. The influence of temperature on the mechanical behaviour was examined. TEM

studies were carried out post testing on typical samples, and the possible mechanisms of DSA

were discussed.

2. Mechanical Testing

Mechanical testing was carried out on an Instron servo-hydraulic testing machine with

computer-controlled loading spectra programmed as required. The temperature was controlled

using a three-zone electrical resistance furnace. The strain was monitored using a high

temperature extensometer (12.5+2.5/-1.25mm). The tests were carried out on smooth

cylindrical specimens with a gauge length of 12.5mm and a diameter of 6.95mm. The

microstructure of CGRR1000 consists of a bimodal γ’ precipitates with cuboidal γ’ precipitates of

a mean size of 280nm and spherical γ’ precipitates of a mean size of 25nm and an overall very

low dislocation density [12, 13]. Full material description and experimental procedure are given

elsewhere [12] and the chemical composition is given in Table 1 [14].

Twelve tests were carried out under uniaxial loading, including strain-controlled multiple

hardening relaxation tests (MHRT), strain rate controlled cyclic tests and simple cyclic tests

under load control. The first series, MHRT, included stepped monotonic loading up to a strain of

2% at a strain rate of 0.5%/s, with a hold period of 100s imposed at every 0.2%strain from 1% till

2% of strain. The tests were performed at room temperature, 550°C, 600°C and 800°C. The

second series of tests were strain rate controlled tests of three segments, from zero to 1% strain

at a strain rate of 0.5%/s; from 1% to 2% at a strain rate of 0.005%/s and from 2% to 3% at a

strain rate of 0.5%/s. This series of tests were designed to have two orders of difference in strain

rate to examine the effects of strain rate on the stress-strain behaviour, which may help to

discern DSA behaviour [15]. This series of tests was carried out at 600°C, 650°C, 700°C, 750°C

and 775°C. The last series of tests were performed cyclically using a fully balanced strain range

of 2% at 300°C and 650°C, respectively. A summary of the test conditions is given in Table 2.

3. Results and discussion

3.1. MHRT series

Figure 1 shows the stress-strain responses from the MHRT series. The aim of the tests was to

evaluate the viscoplastic behaviour of the material by holding the strain constant at selected

strain values and allowing stress relaxation. At room temperature, small but notable stress

relaxation may be observed. At 800°C, the stress relaxation behaviour was much more

pronounced with an average stress drop of more than 200MPa after 100 seconds; and the

behaviour became more evident with increasing strain. At the intermediate temperatures 550°C

and 600°C, however, a distinctly different behaviour was observed. Firstly, the material showed

no gradual stress relaxation as in 800°C, rather jumped up and down to a more or less constant

stress level during the hold time. Secondly, the stress level “settled” at each strain hold seems to

be rather erratic, with lower stress level at 1.2% than that at 1%, for example. The stress-strain

curves at these two temperatures are remarkably similar, a phenomenon termed as Portevin Le-

Chatelier effect [1] due to DSA. This phenomenon is clearly temperature dependent such that

only results from the intermediate temperatures are affected whilst room temperature and

800°C seem to be outside of the temperature range when DSA is active for this material. As

stress relaxation requires sufficient mobility of the dislocations, it is plausible that the absence of

the continuous decrease of stresses observed at the intermediate temperatures is a

consequence of the dislocation pinning due to DSA. In fact, Avalos et al. [16] suggested that the

internal stress from the mechanism associated with DSA may be responsible.

3.2. Strain rate controlled tests

Results from the strain rate controlled test series are shown in Figure 2. It appears that the

general behaviour of the material does not vary much with temperature up to 750°C. The

material seems to be insensitive to strain rate with very little variation in the stress response

even when the strain rate was changed by two orders of magnitude. This is somewhat unusual

as DSA is usually associated with negative strain rate sensitivity [6], hence a significant reduction

in strain rate should result in an appreciable increase in flow stress in the temperature regime of

DSA. According to Brechet and Estrin [17, 18], other mechanisms such as “pseudo PLC effects”

might be at work as well. Moreover, the overall strength of the material is controlled by

homogeneously distributed fine precipitates, which strongly interact with dislocations. These

interaction mechanisms are also likely to mask the DSA phenomenon.

At test temperature 775°C, a significant drop in the stress-strain curve may be observed

associated with the slowest strain rate, suggesting a positive strain rate sensitivity. It is

interesting to note that once the test was resumed with a strain rate of 0.5%/s, the stress

jumped back to the same level as it was before at 0.5%/s, albeit with an overshoot.

A close look at the individual curves shown in Figure 2 reveals a frequent presence of

serrations during loading, mostly associated with high strain rates and at high strains. At 600°C,

evident serrations may be observed as early as 0.6% strain. They disappear during the slow rate

phase but reappear as soon as the strain rate is increased. These serrations may be associated

with acoustic emissions during the two faster loading rates. For the test at 650°C, the early

phase does not show any serrations but they appear as soon as the fast strain rate is resumed

until about 2.2% strain, above which they disappear again. Interestingly, the test at 700°C

displays two types of serrations in the third phase of the test. The first type of serrations, Type A,

of small amplitude, appears early and persists throughout the test. The second type of

serrations, Type B, is of larger amplitude and appears at about the same strain level as at 650°C.

The test at 750°C shows only Type A serrations but they appear early at 0.6% strain and are

present during loading at both 0.5%/s and 0.005%/s. These phenomena seem to be consistent

with those reported in [4, 19]. For Alloy CGRR1000, it seems that Type B serrations, known as

locking serrations, are predominant in the tests at 600, 650 and 700°C;whilst Type A serrations

are found in the tests at 750 and 775°C.

3.3. Cyclic behaviour

Results from the cyclic tests at 300°C and 650°C under a strain range of 2% (fully reversed)

and a strain rate of 0.005%/s are presented in Figure 3. It appears that, at 300°C, the material

shows predominantly cyclic hardening behaviour, albeit at a reducing rate with the number of

cycle, and does not appear to reach saturation even after 50 cycles. This behaviour differs from

that at 650°C , where, after the initial period of hardening, the material stabilises after about 20

cycles with very slight softening thereafter. Some interesting surface features are presented in

Figure 4. For the sample tested at 300°C (Figure 4a), slip lines are clearly visible with lengths

typically between 50 to 80 μm intersecting at about 60° the loading axis. These lines are mainly

localised and appear to be the result of the cyclic loading. In Figure 4b, the sample tested at

650°C shows multiple small cracks as well as slip lines similar to those in Figure 4a. Oxide scales

formed at 650°C might have reduced the overall visibility of the slip lines, although they are still

visible and similar in appearance to those found at 300°C. These slip lines seem to be a

characteristic of DSA, which is usually associated with the localisation of plastic deformation due

to repeated initiation and propagation of slip bands [20].

3.4. Hardening behaviour

The normalised ultimate tensile stress (UTS), calculated by UTS/UTS20°C[21], is shown as a

function of temperature in Figure 5. From room temperature up toabout 300°C, the behaviour

seems to be common to most metals, namely a decrease of UTS with the increase of

temperature. However after 300°C, the UTS climbs up again showing that the material strength

increases with temperature until the temperature reaches about 650°C then the strength

plummets again at higher temperatures. Such a behaviour may well be associated with DSA

taking place in the temperature range between 300°C and 650°C [22, 23].

Further cyclic tests were carried out at a strain range of 2% and a stress ratio of R=-1 and at a

strain range of 3% and a stress ratio of R=0to evaluate the strain hardening behaviour. The

results are shown in Figure 6, where the hardening ratio, defined as ( )/

[19], is plotted as a function of temperature. Results for a full range of temperatures could not

be obtained partly due to the lack of data, partly due to the absence of saturation (as shown in

Figure 3). A peak in the hardening ratio at 650°C in both series of tests seems to be indicative of

the influence of DSA .

3.5. Dislocation arrangements

For the tests of series 1, MHRT, and series 2, cyclic test at 300°C, a transmission electron

microscopy (TEM) analysis of the microstructure has been carried out. In contrast to the results

of the mechanical tests, where the effects of DSA can be clearly observed, TEM results seem to

be less conclusive and these will be discussed below.

At temperaures below 300°C planar dislocation arrangements seem to be dominant, as seen

in Figure 7a for the sample tested at room temperature (MHRT) and Figure 7e for the specimen

tested cylically at 300°C. Between 500°C and 750°C, where the PLC effect is at work and the

maximum strength and hardening achieved at about 650°C (Figures 2, 5 and 6), the dislocation

arrangements are quite different from those at room temperature. Multiple slips are observed

in the γ matrix between the bimodal γ’ precipitates in the specimen tested at 600°C MHRT

(Figure 7c). A similar dislocation arrangement may be observed for the sample tested cyclically

at 650°C, where multiple slips seem to prevaile in the γ matrix, as shown in Figure 7f.

Interestingly, the dislocation arrangement of the sample tested at 550°C MHRT indicates a

transition between the planar to multiple slips, as shown in Figure 7b, where a mixture of planar

dislocation arrangements, similar to those found for the low-temperature tests, and multiple slip

dislocation networks in the γ matrix, as seen in samples tested at higher temperatures.

Furthermore, for the 650°C cyclic test, evidence of dislocation climb seems to be visible from the

TEM analysis. At 800°C, the TEM micrograph (Figure 7d) of the sample tested under MHRT

shows pronounced dislocation climb, indicating high mobility of the dislocations; also

homogenously spread dislocation movement, where precipitates are no longer significant

obstacles, which might otherwise prevent smooth dislocation motion by, for example, pinning.

Such a “smoooth” dislocation movement enabled the observed stress relaxation at 800°C (Figure

1) .

As similar dislcoation arrangements are seen at the same temperature for both series of

tests, dislocation arrangements appear to be independent of testing method. The characteristics

of the dislocation arrangement with temperature are rather typical, i.e., planar dislocation

arrangements at low temperatures, multiple slips at higher temperatures and dislocation climb

at higher still temperatures. The observed dislocation arrangements do not appear to correlate

with the DSA observed during mechanical testing, or known typical DSA dislocation

arrangements for precipitation free materials, such as planar dislocation arrangements or

dislocation loops in 316LN steel [24] or uncondensed and poorly defined walls for DSS [25]. It

might be that the effects of precipitates on dislocations glide mask the visible effects of the DSA

on dislocation arrangements. Indeed, similar dislocation arrangements as in this study were

reported in a precipitation-hardened Ni-Co-base superalloy [26]. Nevertheless, the TEM results

might indicate a pseudo PLC effect, which could be the reason for serrated flows and negative

strain rate sensitivity [17, 18, 26, 27]. The pseudo PLC effect may be explained by either the

shearing of precipitates [17-19] or the deformation induced dissolution of precipitates [26, 27].

Shearing of precipitates has not been observed in this TEM analysis; whilst dislocation climb

dominates in the sample cyclically tested at 650°C and climbing and shearing do not normally

appear simultaneously in a material. Hence deformation-induced dissolution of precipitates as

reason for the pseudo PLC effect is unlikely. In case of a deformation-induced dissolution of the

γ’ precipitates a serrated flow should also be visible in the 20°C MHRT, which is not the case as

shown in figure 1. Therefore, the lack of serrated flow in the 20°C MHRT and the missing

shearing of precipitates lead to the assumption that DSA is more likely as the pseudo PLC effect.

4. Concluding remarks

Dynamic strain aging (DSA), presented as serrated yielding in stress-strain responses,has

been observed in Alloy CGRR1000 under both monotonic and cyclic loading conditions. DSA

mechanism seems to be active in a temperature range from 300°C to 750°C for this alloy,

although it appears to be largelyinsensitive to strain rate within the range of 0.005m/s to

0.5m/s.Contrary to the usual behaviour that serrated flow becomes more pronounced as

temperature rises and strain rate decreases, the serration activity in the current material

appears to decrease with decreasing strain rate. Other mechanisms such as “pseudo PLC effects”

might be at work as well but are unlikely given the TEM results. However, the main reason for

the difference in the DSA responses from CGRR1000 and ferritic and austenitic steels may be the

very strong effect of particle strengthening in CGRR1000 Alloy, which may have masked the

influence of DSA on the dislocation motion and arrangement.

The material behaviour in both multiple hardening relaxation and cyclic tests manifests itself

in dislocation arrangements, which is closely related to the test temperature but largely

independent of the test mode. At room temperature, planar dislocation arrangements prevail.

With increasing temperature the dislocation arrangement becomes increasingly more typical of

wavy dislocation slip, as a consequence of the increased contribution of thermally activated

mechanisms of dislocation motion. The corresponding change in the dislocation arrangement

seems to compensate and somehow mask the effect of DSA on the microstructure, i.e. an

increase in planarity as a consequence of reduced dislocation mobility in the matrix.

Acknowledgments

The work was supported by British Council and DeutscherAkademischerAuslandsdienst and

by the Technology Strategy Board of the UK.

References:

[1] A.Portevin, F.Le-Chatelier, Sur un phénomèneobservélors de l’essai de traction d’alliages

en cours de transformation, ComptesRendus de l’Academie des Sciences, Paris176(1923) 507–

510.

[2] S. Zhang, P. G. McCormick, Y. Estrin, The morphology of Portevin–Le Chatelier bands:

finite element simulation for Al–Mg–Si, Acta Mater. 49 (2001) 1087-1094.

[3]P.G. McCormick, A model for the Portevin–Le Chatelier effect in substitutional alloys,

ActaMetall.20(1972) 351.

[4] P. Rodríguez, Serrated plastic flow, Bull. Mater. Sci.6(1984) 653.

[5] V.S. Srinivasan, R. Sandhya, K.BhanuSankaraRao, S.L. Mannan, K.S. Raghavan, Effect of

temperature on the low cycle fatigue behaviour of nitrogen alloyed type 316L stainless steel, Int.

J. Fatigue 13(1991) 471–478.

[6] L. Fournier, D. Delafosse, T. Magnin, Oxidation induced intergranular cracking and

Portevin-Le Chatelier effect in nickel base superalloy 718, Mater. Sci. Eng. A 316(2001) 166–173.

[7] M. Mazière, J. Besson, S. Forest, B. Tanguy, H. Chalons, F. Vogel,Numerical aspects in the

finite element simulation of the Portevin–Le Chatelier effect, Comput. Methods Appl. Mech.

Eng.199(2010) 734-754.

[8] K. Gopinath, A.K. Gogia, S.V. Kamat, R. Balamuralikrishnan, U. Ramamurty, Low cycle

fatigue behaviour of a low interstitial Ni-base superalloy, Acta Mater. 57 (2009) 3450-3459.

[9] M. Valsan, D.H. Sastry, K.BhanusankaRao, S.L. Mannan, Effect of strain rate on the high-

temperature low-cycle fatigue properties of a Nimonic PE-16 superalloy, Metal. Mater. Trans. A

25 (1994) 159–171.

[10] R.W. Hayes, W.C. Hayes, On the mechanism of delayed discontinuous plastic flow in an

age-hardened nickel alloy, Acta Metall. Mater. 30 (1982) 1295.

[11] R.W. Hayes, On a proposed theory for the disappearance of serrated flow in f.c.c. Ni

alloys, Acta Metall. Mater. 31 (1983) 365.

[12] Cornet, C. Modelling damage and deformation bahviour of nickel-based superalloy using

constitutive models. Ph.D thesis, University of Portsmouth; 2009.

[13] C. Stöcker, M. Zimmermann, H.-J. Christ, Z.-L. Zhan, C. Cornet, L.G. Zhao, M.C. Hardy, J.

Tong, Microstructural characterisation and constitutive behaviour of alloy RR1000 under fatigue

and creep-fatigue loading condition, Mater. Sci. Eng. A 518(2009) 27-34.

[14] S. Everitt, M.J. Starink, H.T. Pang, I.M: Wilcock, M.B. Henderson, P.A.S. Reed, A

comparison of high temperature fatigue crack propagation in various sub-solvus heat treated

turbine disc alloys, Mater. Sci. Techn. 23 (2007) 1419-1423.

[15] M. Yaguchi and Y Takahashi, Unified inelastic constitutive model for modified 9Cr-1Mo

steel incorporatingdynamic strain aging effect, JSME Int. J. A, 42 (1999) 1–10.

[16] M. Avalos, I. Alvarez-Armas, A.F. Armas, Dynamic strain aging effects on low-cycle

fatigue of AISI 430F, Mater. Sci. Eng. A, 513-514 (2009) 1-7.

[17] Y. Brechet, Y. Estrin, On a Pseudo-Portevin-Le Chateliereffect, Scr. Metall. Mater. 31

(1994) 185.

[18] Y. Brechet, Y. Estrin, Pseudo-Portevin-Le Chatelier effect in ordered alloys, Scr. Mater.

35 (1996), 217.

[19] E. Pink, Grinberg A. Leoban, D.F. Mexico, Practical aspects of the Portevin-Le Chatelier

effect, Aluminium 60 (1984) E601–E604.

[20] L.P. Kubin, Y. Estrin, The Portevin-Le Chatelier effect in deformation with constant stress

rate, Acta Metall. Mater. 33 (1985) 397.

[21] Rolls-Royce plc, private communication (2010).

[22] M.R. Akbarpour and A. Ekrami, Effect of temperature on flow and work hardening

behavior of high bainite dual phase (HBDP) steels, Mater. Sci. Eng. A 475(2008) 293-298.

[23] A. Gironès, L. Llanes, M. Anglada, A. Mateo Dynamic strain ageing effects on

superduplex stainless steels at intermediate temperatures, Mater. Sci. Eng. A 367 (2004) 322-

328.

[24] S.L. Mannan, Role of dynamic strain ageing in low cycle fatigue, Bull. Mater. Sci. 16

(1993) 561-582.

[25] S. Hereñú, I. Alvarez-Armas, A.F. Armas, The influence of dynamic strain aging on the

low cycle fatigue of duplex stainless steel, Scr. Mater. 45 (2001)739-745.

[26] C.Y. Cui, Y.F. Gu, Y. Yuan, H. Harada, Dynamic strain ageing in a new Ni-Co base

superalloy, Scr. Mater. 64 (2001) 502-505.

[27] K. Gopinath, A.K. Gogia, S. Kamat, U. Ramamurty, Dynamic strain ageing in Ni-base

superalloy 720Li, Acta Mater. 57 (2009) 1243-1253

Table 1: Composition of CGRR1000 [14]

Element Co Cr Mo Al Ti Hf C B Fe Zr Ta

(wt. %) 14.0-19.0

14.35-15.15

4.25-5.25

2.85-3.15

3.45-4.15

0.5-1.0

0.012-0.033

0.01-0.025

0.0-1.0

0.05-0.07

1.35-2.15

Table 2: Summary of the Tests

Test typeTemperature

(°C)Sample name

Cyclic (2%, R=-1)

300 HT010

650 HT011

650 HT014

Strain rate controlled

(0.5-0.005-0.5%s-1)

600 HT017

650 HT012

700 HT015

750 HT016

775 HT013

MHRT

(100s dwell at

1-1.2-1.4-1.6-1.8-2%)

23 HT008

550 HT007

600 HT001

800 HT003

Figure 1: Multiple hardening relaxation tests (MHRT) at selected temperatures.

(a) (b)

(c) (d)

(e)

Figure 2: The individual stress-strain responses of the strain rate controlled tests at

(a) 600; (b) 650; (c) 700; (d) 750 and (e) 775°C.

Figure 3: The stress range evolution with the number of cycles at 300°C and at 650°C.

Strain range 2%, R=-1; strain rate 0.005%/s.

(a) (b)

Figure 4: Micrographs of the sample surfaces post cyclic testing (series 3)

at (a) 300°C and (b) 650°C.

Figure 5: Evolution of the ultimate tensile stress (UTS) as a function of temperature [19].

Figure 6: The hardening ratioas a function of temperature.

(a) (b)

(c) 250 nm250 nm 250 nm250 nm (d)

(e) (f)

Figure 7: Dislocation arrangements: (a) Planar dislocation arrangement for the room temperature MHRT, (b) planar dislocation arrangement and multiple slip for the 550°C MHRT, (c) multiple slip for the 600°C MHRT, (d) dislocation climb for the 800°C MHRT, (e) planar dislocation arrangement for a cyclic fatigue test peformed with balanced strain of 2% at a strain rate of 0.005%/s at 300°C and (f) multiple slip concentrated in the matrix between the precipitates for the a fatigue test performed with balanced strain of 2% at a strain rate of 0.005%/s at 650°C.

Highlights:

Dyamic strain aging (DSA) has been observed in a nickel base alloy CGRR1000 at temperature range 300-750°C.

DSA appears to be insensitive to strain rate and independent of the loading mode for the material and test conditions examined.

Dislocation arrangements appear to be strongly associated with the presence of precipitates and are less affected by DSA.