Resisting stress of a low alloy ferritic steel after creep exposure in service

Acta metall, mater. Vol. 40, No. 10, pp. 2607-2616, 1992 0956-7151/92 $5.00 + 0.00 Printed in Great Britain. All rights reserved Copyright ~, 1992 Pergamon Press Ltd

RESISTING STRESS OF A LOW ALLOY FERRITIC STEEL AFTER CREEP EXPOSURE IN SERVICE

R. SINGH and S. BANERJEE Materials Evaluation Division, National Metallurgical Laboratory, Jamshedpur 831 007. India

(Received 25 September 1991; in revised form 21 February 1992)

Abstract--Creep data generated on the materials in the service exposed and the virgin states of a 1 CRY.3 Mo4).25 V steel and also on materials with a wide range of microstructures prepared from the virgin state, have been analysed to determine the resisting stress. The analysis shows that for a given microstructure, as the applied stress increases and reaches a critical value called the transition stress, the creep rate increases sharply from a region of low stress sensitivity to a region of high stress sensitivity. Correspondingly, the resisting stress which increases with applied stress, reaches a threshold value. The results show that, as expected, the transition stress and also the resisting stress in the high stress sensitivity region, depend on the microstructure. However, a remarkable observation is that the resisting stress in the low stress sensitivity region is practically independent of the initial microstructural states investigated. Also the threshold stress is observed to be about 0.5 times the hot yield strength. Based on these observations, a simple method for creep life prediction is proposed.

R6sum6--On analyse, afin de d&erminer la contrainte de r6sistance, des donn6es de fluage relatives 5. un acier 1Cr-0,3 Mo-0,25 V en service ou vierge, ainsi qu' 5, des mat6riaux pr6sentant une large gamme de microstructures et pr6par6s 5, partir de l'6tat vierge. L'analyse montre que, pour une microstructure donn6e, lorsque la contrainte appliqu6e croit et atteint une valeur critique appel6e contrainte de transition, la vitesse de fluage croft brutalement depuis une r6gion de faible sensibilit6 5. la contrainte jusqu'fi une r6gion de forte sensibilit6. En m~me temps, la contrainte de r6sistance, qui croft avec la contrainte appliqu6e, atteint un seuil. Les r6sultats montrent que, comme pr6vu, la contrainte de transition et la contrainte de resistance dans le domaine de forte sensibilit6 fi la contrainte, d6pendent de la microstructure. Cependant, il 6st important de noter que la contrainte de resistance dans le domaine des faibles sensibilit6s 5, la contrainte, est pratiquement ind6pendante des 6tats microstructuraux initiaux 6tudi6s. On observe aussi que la contrainte de seuil est environ 0,5 fois la limite 61astique 5, chaud. En se basant sur ces observations, on propose une m&hode simple pour pr6voir la dur6e de vie en fluage.

Zusammenfassung--Kriechdaten des Stahles ICr~), 3Mo4),25 V, im Beanspruchungs- und Neuzustand und mit verschiedenen, am Neuzustand pr~parierten Mikrostrukturen, werden analysiert, um den Verformungswiderstand zu bestimmen. Die Analyse zeigt, dab die Kriechrate bei einer gegebenen Mikrostrucktur yon einem Bereich niedriger Spannungsabh~ingigkeit scharf in einen Bereich hoher Spannungsempfindlichkeit fibergeht, wenn die ~iuBere Spannung ansteigt und einen kritischen Wert, die Obergangsspannung, erreicht. Entsprechend erreicht der Verformungswiderstand, der mit der ~iuBeren Spannnng zunimmt, einen Schwellwert. Die Ergebnisse zeigen, dab t3bergangsspannung und Verfor- mungswiderstand im Bereich hoher Spannungsabh/ingigkeit wie erwartet vonder Mikrostruktur ab- h~ingen. Allerdings ist die Beobachtung bermerkenswert, daB der Verformungswiderstand im Bereich niedriger Spannungsabhiingigkeit vom Anfangszustand der untersuchten Mikrostrukturen praktisch nicht abhfingt. AuBerdem ergibt sich, dab die Schwellspannung etwa 0,5 der Warmstreckgrenze ist. Auf der Grundlage dieser Beobachtungen wird eine einfache Methode zur Bestimmung der Kriechstandfestigkeit vorgeschlagen.

1. INTRODUCTION

The concept of resisting stress, pioneered by Wilshire and co-workers [l, 2], has been widely used in the literature to rationalize the creep behaviour of particle strengthened alloys. For example, a large body of creep data on C r - M o - V steels [3, 4] and Ni-base alloys [5-8] have been fitted to a semi-empirical power law creep equat ion modified by using the concept of resisting stress. This is reproduced below

AD#b ,, ~m = ~ t t o " -- trR)/kt] "° ( l)

where ~m is the minimum creep rate, a R is the resisting stress, a is the applied stress,/a the shear modulus, b is the Burgers vector, k is the Boltzmann's constant, T is the absolute temperature, D is the bulk self diffusion coefficient, A is a dimensionless constant and n o is the modified stress exponent which is equal to about 4.

The resisting stress, tr R has also been referred to as back stress [8] or friction stress [5]. In this paper, the term resisting stress will be used subsequently.

The term resisting stress and the other related parameters such as the transition stress, ax and the

2607

2608 SINGH and BANERJEE: RESISTING STRESS OF A LOW ALLOY STEEL

threshold stress, a H can be conveniently understood from the schematics reported in Fig. l(a,b).

Figure l(a) represents a large volume of creep data [3-7, 9] in a log-log plot of the minimum creep rate vs the applied stress. The figure shows that as the applied stress increases to a value t~ = aT, the minimum creep rate increases very sharply from a region of low to a region of high stress sensitivity, where t~ r is called the transition stress. It is widely believed that, at this stress, the mechanism by which a dislocation overcomes precipitates undergoes a transition from one of predominantly climb to a predominantly bowing type [5, 6, 9].

Figure l(b) shows a plot of experimentally determined values of aR VS the applied stress [9-11]. The figure shows that for a < aT, that is in the region of low stress sensitivity, aR increases linearly with a. However, at a > aT, erR is a constant which is called the threshold stress, t~ H . This is called threshold stress since at a < all, very high creep rates cannot occur. On the other hand, at a >aH (or more appropriately at a > aT) creep rate increases rapidly producing the high stress sensitivity of the creep rate. The different procedures used for the determination of aR are compared later.

William and Wilshire [12] had explained that the early onset of tertiary creep or very high rates of creep and its dominance during the service exposure of an engineering component leading to the creep failure, is caused by a steady decline in the resisting stress as a result of the coarsening of carbides in ferritic steel. However, a recent study [13, 14] shows that even after 10 years of creep exposure at 540°C/50 MPa, the change in composition and morphology of the carbides in a 1 Cr--0.3 Mo-0.25 V grade of steel is not significant. This is further confirmed by the obser-

E ° ~

0 - J

n-,.-16-40

I i

O T

Log o

{b) OR = KO-.~----- I D o R = o H

( OH '

o T o

Fig. 1. Schematic representation: (a) the dependence of minimum creep rate on applied stress [3-7, 9] and (b) the

dependence of resisting stress on applied stress [9-11].

vation that the rupture strengths of the virgin and service exposed samples at low stress levels are almost same [13, 14]. Thus, it seems that the microstructural changes caused due to about l0 years of service exposure of an engineering component, does not produce significant creep degradation. Therefore, the creep behaviour of the ferritic steel in the regime of stress and temperature actually experienced by the engineering components in a power plant, appears to be different from what is usually assumed.

The trend of this unexpected behaviour should be further confirmed by investigations on specimens prepared from components exposed to service for 20 or even 30 years. However, due to the obvious difficulties, such studies have hardly been carried out. On the other hand, the behaviour described above can be checked, even though indirectly, by analysing the creep data generated on specimens with different microstructural states. This enables one to determine the effect of the initial microstructural state on the resisting stress.

Despite its significant use in representing creep behaviour, some important questions relating to the resisting stress concept remain unanswered. For instance: is the concept applicable to the Cr -Mo-V ferritic steel, or in other words, do such steels have a clearly identifiable a T (or all) above which e~n increases sharply? How does the long term service exposure or a change in the initial microstructural state, influence cr R in the low and in the high stress sensitivity region? How does a change in microstructure influence the values of a T and tTn? Can o" H value represent the extent of microstructural degradation? Can aH be correlated to the high temperature yield strength? The main objective of this paper is to examine these questions.

Accordingly, this paper reports the results of the analysis of creep data generated on the materials in the virgin and the service exposed states and also on the materials in the virgin state transformed into three widely different microstructures with significantly different tensile properties. Both the service exposed and the virgin material belong to the same manufacturing batch of a l Cr-0.3 Mo-0.25 V ferritic steel. In addition, creep data reported recently for a 2.25 Cr-1 Mo steel [15] have been analysed to com- pare the same with the trend of results reported here for the 1 Cr-0.30 Mo-0.25 V steel.

2. EXPERIMENTAL

Material

The experimental material is the main steam pipe of a 1 Cr-0.3 Mo-0.25 V steel (Russian grade 12 kh 1 MF) received from an Indian thermal power plant both in the virgin (VN) and the service exposed (SE) states. The samples were cut from the pipe in the VN state, and in the SE state, that is after 105 h (about 10 years) of service exposure at 540°C under a hoop stress of 50 MPa. Both the VN and the SE pipes

S I N G H and B A N E R J E E : R E S I S T I N G STRESS O F A L O W A L L O Y S T E E L

Table 1. Details of the lCr-O.3Mo-0.25 steel samples investigated

2609

State/ Heat Hardness Sample dimension treatement Microstructure (HV)

VN Virgin pipe: Commercially 90% ferrite + 156 273 mm OD x 22 mm thick N at 980 ~ + T at 700°C 10% bainite

SE Same as VN but service - - 90% ferrite+ 152 exposed at 540~C/50 MPa 10% tempered bainite (hoop stress)

RHVI 15 mm dia bars prepared Austenitized at 980 'C (1 h) 100% baimte 210 from the VN material followed by oil quenching

and T at 700':C (2 h) RHV2 As RHV1 N at 980':C (1 h)+ 80% ferrite+ 190

T at 700"C (2 h) 20% bainite RHV3 As RHV1 N at 980:'C (1 h)+ 80% ferrite+ 160

T at 700~C (72 h) 20% tempered bainite

N normalized: T tempered.

belong to the same manufacturing batch. The extensive metallography carried out on these samples have been reported earlier [13, 14]. The SE samples did not contain any intergranular cavity.

In addition to the VN and the SE samples, three more sets of samples having three different microstructures were prepared in the laboratory by suitable heat treatments of the VN samples. The details of the treatment, microstructure and the initial hardness of all the five samples are listed in Table 1. The chemical composition of the steel is given in Table 2.

Table 1 shows that the five samples represent five different microstructural states--varying from predominantly ferritic to predominantly tempered bainitic structure.

Creep-rupture test

Uniaxial creep tests were conducted on 8 mm diameter × 50 mm gauge length specimens machined with their length oriented along the longitudinal axis of the steam pipe. The creep tests were conducted at 550'~C in air and under constant load at stress levels varying between 69-294 MPa in creep testing ma- chines having automatic lever levelling facility. Creep strains were measured by LVDTs with strain resolution of 0.01% and the test temperature was con- trolled within +2°C.

High temperature tensile tests

High temperature tensile tests were conducted in a screw driven universal testing machine with strain resolution and temperature control equivalent to that used in the creep testing. Tests were carried out on

Table 2. Chemical composition (wt %) of the pipe investigated

C Mn P S Si Cr. Mo V N

0.11 0.30 0.01 0.007 0.35 1.01 0.28 0.23 0.023

Table 3. Tensile properties at room and elevated temperatures for 1Crq).3Mc~0.25V steel

0.2% proof Elongation Reduction Temperature stress U.T.S. (on 50 ram) in area

Sample (~C) (MPa) (MPa) (%) (%)

VN Room 360 480 25 74 350 240 476 30 79 450 226 432 30 79 500 210 323 32 78 550 206 304 30 75 600 165 232 35 79 650 128 186 53 94

SE Room 280 475 26 72 300 206 383 22 75 400 216 383 24 68 500 206 304 34 71 550 177 255 36 76 600 157 215 48 83 650 130 183 55 89

RHVI Room 667 794 17 70 300 598 736 19 72 400 549 686 18 73 450 530 637 19 74 500 500 588 18 79 550 461 520 22 83 600 363 399 26 90

RHV2 550 337 445 17 83 RHV3 550 231 373 18 83


duplicate specimens in the five different microstructural states at a constant strain rate of 10 -3 per s. Table 3 reports the high temperature tensile test properties.

3. RESULTS AND DISCUSSION

The data obtained are analysed to evaluate tra and aT values, in the five different microstructural states. The possible correlations between tr R and tr and that between tr H and the yield stress at the creep testing temperature, are examined and discussed. The trend of results observed here in case of the 1 Cr--0.3 Mo--0.25 V, has been compared with that for the data recently reported on 2.25 Cr - I Mo steel [15]. Finally, based on the correlations established, a very simple procedure of life prediction is proposed.

Minimum and steady state creep rates

Figure 2 reports the creep data at 550°C for the various microstructures in a manner similar to that reported in the schematic in Fig. l(a).

The minimum creep rate data used in Fig. 2 needs some clarification, since a distinction has been made between the minimum and the steady state creep rate in the literature [16-20]. For the creep strain values up to 1%, which is the maximum in-service strain permitted by the design codes of a power plant, the creep curves of all the five microstructures studied here, exhibit a behaviour which is termed as "concave downward" [20] behaviour. This means that upto 1% strain, the steady state creep continued, the tertiary creep did not begin and therefore the microstructure can be assumed to be stable [20]. A similar behaviour is also exhibited by the data of 2.25 Cr-1 Mo steel recently reported [15] and examined by us. In such an

Appl ied stress, o (MPa) 63 100 150 200 250 300

I I I I ] 550°C x SE

- 3 - - o VN

IIIW 1 n = 1 6 1 6 1 6 16 J ' 5

, ..=v= /g ! / t - x I ¢

-7 ~ = ' Climb Orowan

~limb t, Orowan I

I', I I I ] I I B8 .8 1.9 2.0 2.1 2.2 2.3 2.4 2.5

Log appl ied stress, o (MPa)

Fig. 2. Transition stress for different microstructures.

I t -

v - 5

.=E t ~ O

. - I - 6

instance, the minimum and the steady state creep rates are equal. Therefore, a distinction between the two creep rates of a steel may not be significant if one uses creep rate to predict the life of a component in a power plant.

In this paper, the steady state creep rate data have been used to plot Fig. 2.

a r values for the different microstructural states

The value of apparent creep index [8], n, determined from the slope of the plots in Fig. 2, increases abruptly from about 4 at a < tr r to a significantly high value, about 15 to 16 at a > aT. This behaviour has also been observed in the case of precipitation hardened alloys [5-7] including 1 Cr--0.5 Mo-0.25 V ferritic steel [3, 4].

Such an increase in the value of n from the region of low to the region of high stress sensitivity (see Fig. 2), is generally attributed by most of the re- searchers to a change in the creep deformation mechanism [5,6,21] from climb to Orowan bowing. However, it should be noted that experimental evidence is not completely conclusive since there is a controversy as to whether the Orowan loop, in a sample crept under applied stress exceeding trr, is actually present [5] or not [22].

It is noteworthy that at tr > trT, n increases both in the virgin state (VN) as also after a service exposure (SE) for 105h. The aT has decreased--but quite unexpectedly, only by a small amount (12%) after 105 h of service exposure--confirming that the microstructural changes caused by the service exposure, has produced only a minor amount of creep degradation [13, 14].

This is further confirmed by the fact that the rupture strength data generated on the VN and the SE samples exhibit the so-called "equi-strength state" [13, 14]. Equi-strength state [13, 14] means that the creep rupture strengths at 550°C/50 MPa, as obtained for the VN and the SE samples of the steel investigated in this study, converge. Such convergence of the data is also reported in other ferritic steels, for example in 0.5Cr--0.5 Mo--0.25V [23,24] and 2.25 Cr-1 Mo [25] steels. The convergence confirms that the microstructural degradation in ferritic steels would not produce any appreciable acceleration in the steady state creep rate at the usual 550°C/50 MPa service exposure which is experienced by several critical components in a power plant, such as the main steam pipe, headers, superheaters etc.

The most remarkable feature of the plots in Fig. 2 is that in the low stress sensitivity region, that is at a < a T, the data points for all the different initial micro-structures fall on a single line represented by the apparent creep index value of about 4.

Determination o f a R

The method for determination of resisting stress, tr R, has not been firmly established. The methods more frequently used, such as the stress drop and the

SINGH and BANERJEE: RESISTING STRESS OF A LOW ALLOY STEEL 2611

stress relaxation behaviour have recently been criti- cized [26]. The limitations of these techniques are that they cannot determine (TR in si tu without changing the variables such as stress, temperature or strain rate. In addition, there are other specific limitations, for instance the value of (TR obtained by the stress-drop method depends critically on the sensitivity of the strain measurement [27]; the arbitrary nature of assessing the asymptotic behaviour gives 15% higher value of (7 R [28] in the stress drop technique; and both these techniques require experiments which take long time [29].

Besides these two methods, one can also determine (7 R from the creep data using a method proposed by Lagneborg and Bergman [9]. The method is based on the assumption that as the applied stress, (7, increases, the creep deformation mechanism undergoes a transition from climb to bowing at (7 > (7 T [see Fig. l(a)] and correspondingly. (7 R reaches a maximum or threshold value (7 H. At (7 < (7 T, (TR= K(7 where K = (7 n/(TT.

In the present study, (7 R is determined using Lagneborg and Bergman method. This is shown schematically in Fig. 3, wherein the creep data are plotted in terms of ( g m ) l ;n° VS (7 both in the low (Line 1) and in the high (Line 2) stress sensitivity regions. Equation (1) indicates that such a plot would produce straight lines in both the high and the low stress sensitivity regions. The extension of Line 2 to a strain rate equal to zero gives the value of (7 n .

In the above method, the value of no used is critical. Some authors have preferred n o --= 3.5 [30] and others n o = 4 [311.

This method has also been used recently by Reppich et al. [5] to determine (TR for Ni-base alloys, by Ecob and Evans [22] for dispersion hardened austenitic stainless steel, and by others [30, 32] for dispersion hardened alloys, where n o is assumed to

o

E E t--

E "6

o t,-

O LL

/ Line no. 2 _ /

k i n e ~

/ / ii I

O H

Applied stress o

Fig. 3. Schematic representation: determination of threshold stress from creep data for a given microstructure.

150

125

~'- 100 x

% 75 I t--

X ..,%E 50

25

550°C × SE

- - oVN

- - • RHV 3 ]x /

- / y,

3 X 10 -4 o • ! , , ~ j -

I" i l

,/ I ill 0 50 100

SE o H

B C D E /

~/_ . d

I I , OT

D~-~D I I I I

i I i I

i I i I

I I

I/ i, I

150 200

t RHV1

o H Applied stress, o (MPa)

1 250 300

Fig. 4. Threshold stress for different microstructures in 1Cr-0.30 Mo~).25 V steel.

be 4. The a H values determined by us using this method are discussed subsequently.

a H values ./or the di f ferent mieros t ruc tura l s ta tes

As shown in Fig. 4, a replot of the creep data of Fig. 2 in terms of (ira) °25 VS a produces a linear relationship both in the high and in the low stress sensitivity regions for each of the microstructural states. This shows that Lagenborg and Bergman [9] approach can be used to estimate a H (or at) for a wide range of microstructures produced from the Cr-Mo-V steel.

The intercepts of the lines A, B, C, D and E with the abscissa (Fig. 4) give the values of threshold stress, a n , and the intersection of these lines with the line OM gives the transition stress, aT, for each of the five different microstructures. The lowest a n value of 80 MPa, corresponds to the ferritic microstructure developed in the SE sample and the highest (7H value of 200 MPa corresponds to the bainitic microstructure (RHV1).

It may be noted that the values of (TH determined from Fig. 4, are in general agreement with the values determined using inter-particle spacing by Foldina et al. [23] for the 0.5 CRY.5 Mo4).25 V steel at 550°C. Similarly, the highest ~H value of 200 MPa corresponding to the fully bainitic structure (RHV1) as determined in this investigation, agrees with that determined for the fully bainitic 1 Cr-0.5 Mo~).25 V steel specimen by the stress drop technique [3].

A comparison of the VN and the SE samples in Fig. 4 shows that the value of (7 n decreases only mildly from 92 MPa to 80 MPa after a service exposure of 10 s h. However, if one were to consider that the stress level experienced by a component in a power plant is only about 50 MPa, the effect of such


Table 4. Threshold and the transition stress values as determined from Fig. 4

o- H o- T Specimen (MPa) ( M P a ) aa/a~:

SE 80 95 0.84 VN 92 110 0.83 RHV3 111 130 0.85 RHV2 162 190 0.85 RHVI 205 250 0.82

a decrease in the threshold stress in producing rapid creep rate or creep degradation of the component would not be significant. This contention is also consistent with the observation that the RA% and EL% values (see Table 3) of the VN and SE tensile specimens tested at high temperatures are not significantly different.

It is likely that the intermediate microstructure produced by degradation during service could be appropriately characterized by the an values ranging anywhere between 80 to 200 MPa. Thus, a n of a steel could probably indicate microstructural degradation during long term creep exposure in service. Further work should, however, be carried out with components exposed for 2 x 105 h or more, to establish the validity of this point for microstructures which can have o.n values less than 80 MPa and also for the other ferritic steels.

In terre la t ion o f a n wi th a T

Table 4 reports the values of o. H, o. T and o.~/o. T as obtained from Fig. 4. It is seen that the ratio, o.n/o. T

is about 0.8 for all the microstructural states studied. It is evident from the geometry of the two straight lines reported in Fig. 3, that if Line 1 were to pass through the origin, the ratio of OO.H/O" T will be equal to (m 2 - m l ) / m 2 where ml and m 2 are the slopes of the Lines 1 and 2 respectively. Naturally, the data when plotted in the format as given in Fig. 4, conforms to this requirement.

Interestingly, the values of O'H/aT in Table 4 is in reasonable agreement with those reported in the literature [5, 9, 22] for a range of Ni-base precipitation hardened alloys as well as for stainless steel (Table 5) which were obtained using a procedure identical to the one used in the present study.

The results regarding the relationship between aH and aT are consistent with Ecob and Evan's model [22]: the creep by the Orowan by-pass occurs only at a transition stress which is slightly higher than the

Orowan stress and is high enough to keep 1/3 of the gates [33] open even when they are obstructed by 4-5 Orowan loops accumulated around the particles. These authors have argued that when the applied stress is equal to the Orowan stress (~O'H), about a third of the gates are open initially, but some are then closed by the Orowan loops produced during strain- ing. Thus the Orowan by-pass mechanism ceases when o. ~ o-n.

In the absence of confirmatory evidence (e.g. TEM micrographs), the exact change in the mechanism is not clearly understood. However, the constancy of the an/o.T ratios in different precipitation hardened alloys (Tables 4 and 5) indicates that O'H must be exceeded by a certain amount for the onset of "runaway" creep deformation. Hence, as contended by several authors [34-37], when a > aT, the resisting stress remain constant and is equal to o-H.

Correlat ion o f o. H wi th ho t y ie ld s t rength

aH can indicate the extent of microstructural change during creep exposure. It has also been contended that a H "can be used as a design-limiting strength property" [39]. Notwithstanding this contention, it should be noted that to achieve significantly long life, the stress level in a power plant component used in high temperature service, must be considerably lower than o.H, if long life is to be expected from it in service. On the other hand, O'H represents an upper-bound value of the permissible stress, since the service stress at any point of time must always be less than o. n to avoid runaway creep deformation. Therefore, a simple experimental method to determine o. u should be useful. The metal- lographic methods to determine a H are difficult to carry out and yield uncertain results.

Hence, the correlation of o-n with the hot tensile yield strength as observed by us is significant. This is plotted in Fig. 5 (o-n and hot yield strength values are reported in Tables 4 and 3 respectively) which indicates that

a n = 0.48o. v. (2)

The fact that the plot passes through origin probably has a basic significance, which requires further study.

Similar observation has been made in Ni-base superalloys [8] where o. n = 0.7o.v; the authors [8] have also theoretically predicted the tensile Orowan stress

Table 5. Ratio of a H and o T determined

Alloy system

from the creep data reported in the literature

Test temperature

CC) a H #r x Reference 1. Ni-18.4Cr-2.6AI 10 750 0.75 [9] 2. Ni 19.3Cr-0.83AI-2.06Ti 10 750 0.75 [9 l 3. Ni 19.3C~0.73AI-2.33Ti - - 750 0.75 [9] 4. Nimonic PEI6

(a) Particle dia 36 nm 20 810 0.76 [5] (b) Particle dia 130 nm 20 810 0.82 15]

5. 20Cr-25Ni 1.06Ti 2.96 850 0.70 [22]

Volume fraction of dispersoid

(%)


250

~ - 200

- r

o~ 150 o~ O

~ 100 0 r'-

if/ 0 ¢-

F- 50

550°C x SE o VN zx RHV 1 [] R H V 2 • R H V 3

i /

×

(~H ~ 0 .48 O y

," [ I I I 0 100 200 300 400 500

0 .2% proo f s t ress, Oy (MPa)

Fig. 5. Threshold stress vs hot yield stress at 550 C.

and found that it is nearly equal to o- H. This indicates that the threshold stress is nearly equal to the Orowan stress and scales with the hot yield strength. Also, it can decrease to a lower limit of about one-half of hot yield strength, as temperature increases and/or strain rate decreases [39].

Though the Orowan stress in the present steel cannot be precisely estimated due to a wide variation in the interparticle spacing, an estimate of the upper and the lower bound values should be useful. If the interparticle spacing varies between 100-300 nm, the corresponding tensile Orowan stress (2/~ b ;2) would range between 252 and 84 MPa: where it at 550 C is assumed to be equal to 64 GPa [40]. Interestingly, the lower bound value of 84 MPa, obtained for 2 = 300nm, which is nearly equal to the size of the precipitate free subcells in the VN and the SE samples [13, 14], compares quite well with the experimentally determined c~ H values 80 and 92 MPa (see Table 4).

The high temperature yield stress decreases with decreasing strain rate [41]. Accordingly, a M determined from the steady state creep rate data with a strain rate as low as six orders of magnitude slower than that typically used to determine the hot tensile yield strength, is expected to be significantly lower than ~ry.

Resisting stress and creep rate at a stress less than ar

Figure 6 shows the creep rate data at 550 C for all the microstructures, plotted as (~/m) °25 VS a, in the low stress sensitivity region. These data points are ex- tracted from Fig. 4 and corresponds to the Line 1 of Fig. 3. Two important points should be noted:

1. The data points corresponding to all the five microstructural states, collapse into a single line OM.

2. The best fit lines for the data passes through the origin which can be represented by the equation

(g~m)0 25 = M a (3)

where the experimental value of M determined using least square fit of the data, is reported in Table 6.

The above observations indicate that the creep rate, and therefore ~r R at 550 C, is independent of the initial microstructure as long as a < al-.

As shown schematically in Fig. 1, the resisting and the applied stress are related [9, 10, 41] through

aR = Ko- (4)

where the value of K which is equal to (all /aT), is about 0.8. It is almost the same for all the microstructural states studied (see Fig. 4 and Table 4).

The above result requires some clarification. Be- cause, contrary to the trend of results which shows that the line OM in Fig. 6 passes through the origin, local and general climb theories [16, 21, 42] predict that the plot should show some evidence of an absolute threshold [21].

The local climb has been observed to occur only in oxide dispersion hardened alloys which show increas- ing stress sensitivity with decreasing stress [2l]. As against this, a precipitation-hardened Ni-base superalloy exhibits an opposite trend and the mode of creep deformation in this case is believed to be general climb [21]. The 1Cr4).3 Mo-0 .25V steel exhibits a trend similar to that of precipitation-hardened Ni-base superalloy and one may, therefore, assume that the creep deformation mechanism in this case too occurs by general climb.

The use of general climb theory is known to give two alternative results [9, 21,42]. First, the resisting stress cr R is proportional to cr [9, 21], like in equation (4) above and such a result is consistent with the scheme of life prediction proposed subsequently in this paper. The second result is that the creep

150

125

% I00

x

d ~ 75 e-

x

E 50 ° ~

25

x SE o VN

--z~ RHV 1 cl RHV 2 • RHV 3

M

_ (~m)025 = 3 x 10 -4 o ~ / ~ / / / ~50°C

.

/ ~ I I I [ I 50 1 oo 150 200 250 300

App l i ed st ress, a (MPa)

Fig. 6. Minimum creep rate in the low stress sensitivity region.

AMM 40,10~J

2614 S I N G H and B A N E R J E E : R E S I S T I N G STRESS O F A L O W A L L O Y S T E E L

Table 6. Comparison of experimentally determined and calculated M values of ICr~.3Mo-0.25V steel

M values, [h]°'25/MPa

From experimental data Calculated from equation (5) Microstructure fitted to equation (3) Upper bound a Lower bound b

All five microstructures 3 x 10 -4 7.5 x 10 -4 2.1 x 10 _4

aM (upper bound) is calculated using: ~ = 0.6 [9]; D O = 5.8 x 10 -4 m2/s; Q = 246 kJ/mol [46] and therefore, D at 550°C = 1.8 x 10 -19 m2/s.

bM (lower bound) is calculated using: ~t = 1.2 [9]; D O = 3.4m2/s; Q = 322 kJ/mol [47] and therefore, D at 550°C = 1.7 × 10-2° m2/s.

Assumptions made in the calculation of M from equation (5): resisting stress, tr R = 0.8a x 106 N/m2; shear modulus, p = 64 x 109 N/m 2 [40]; mean free path of dislocation motion, r = 2.5 x lO-Tm [48].

behaviour at very low stress--as represented in Figs 4 or 6, should exhibit an absolute threshold whose magnitude is rather low [43, 44]. In fact, the absolute threshold stress can be as low as 0.03 to 0.07 times aH [16]. It is very difficult to determine such a low threshold stress value experimentally. For instance, with its tr H = 92 MPa, the maximum value of this absolute threshold stress for the VN material would be a mere 6.4MPa. Therefore, even if such a threshold value were to exist, its contribution to the calculated life of a component used in the power plant would be negligible.

An examination of equation (3)

How does equation (3) fit into the available theory? This question must be properly addressed since the observation that the creep rate is practically independent of microstructure in the low stress regime has a significant potential for application in life prediction. Accordingly, the comparison of equation (3) with two well-known constitutive equations for creep rate reported in the literature is made to provide a firmer basis for equation (3).

First, the popular Dorn type equation as given in equation (1) contains a dimensionless constant A. Since A is reported to vary within a very large range, that is between 1-1016 [45], it is very difficult to assign a physical significance to A. However, it may be noted that the creep rate calculated from the Dorn equation agrees with that given by equation (3) if the value ofA = 10. This is in general agreement with the estimated values of A for Cr -Mo-V steels [3, 4].

Second, the Lagneborg and Bergrnan [9] equation

~m = (O#T/2kT)[(tr -- CrR)/C~#]"° (5)

contains a dimensionless constant, ct, which varies within the range 0.6-1.2, depending on the strength of dislocation obstacles [9]. Besides the equation contains, an additional parameter: the mean free path of dislocation motion, [, which has to be estimated to a first approximation. However, equation (3) can be more meaningfully compared with equation (5).

Table 6 reports the upper- and the lower-bound estimates of M values using equation (5) for all the mierostructures investigated. The bases of such esti- mations are described in the footnotes of Table 6.

It may be seen from Table 6 that the value of M obtained from the experimental data fitted to

equation (5) is equal to 3 x 10 -4 and lies within the upper- and the lower-bound M values calculated from equation (5).

An obvious question is, does M change with the type of ferritic steel tested?

Comparison of the behaviour of 1 Cr-O.3 Mo-O.25 V and 2.25Cr-1 Mo steels

The creep rate data reported in this paper is for only one steel--namely 1 Cr-0.3 Mo--0.25V steel. The trend of results reported here for this steel should be confirmed. This is rather difficult since such creep rate data not only for this, but even for the ferritic steels, are rarely reported. Recently, the minimum creep rate data for a 2.25 Cr - I Mo steel have been systematically obtained and reported [15] even though the details of the microstructure is not known.

The 2.25 Cr - I Mo ferritic steel [15] displays a log (7 VS /~m behaviour which is similar to that of 1 Cr-0.3 Mo-0.25 V steel but the increase in n value at a stress above aT is somewhat less, that is n increases from 4 to 7 instead of 4 to 16 in case of the 1 Cr-0.3 Mo-0.25 steel.

150 l

Wolf's data (15): / 125 -- 2.25Cr- 1 Mo/550°C /

~'~ lOO X

(5 - - 75

t -

• 0 . 2 5 x ( ~ )

• 50 - = 6 × I 0 - 4 o

- - I i I I I ~ |

/ 7

/ I / I I I I 0 501~ 100 150 200 250

o H

Applied stress, ~ (MPa) Fig. 7. Thresho ld stress of 2.25 Cr -1 M o steel based on d a t a

repor ted in Ref. [15].


To check the trend of the results reported in Fig. 4, the creep rate data reported for the 2.25 Cr-1 Mo steel [15] have been replotted in Fig. 7. Like in Fig. 4, the replotted data produces two different lines.

Figure 7 shows that the 2.25 Cr-1 Mo steel experi- ences rapid creep above aT= l l 0 M P a and gives a clearly identifiable on value of 55 MPa. The value of an~aT in case of 2.25 Cr-1 Mo, is low: 0.5 instead of 0.8 in the case of 1 Cr-0.3 Mo~.25V. However, a value of OH~aT which is as low as 0.5, is not unexpected [9, 21]. Equation (4) shows that the decrease in O H / a T would correspondingly decrease aR/o from 0.8 to 0.5.

Equation (6) would indicate that with other parameters being equal, a decrease in aRIa [which is equal to K in equation (4)] from 0.8 to 0.5 would increase M value of equation (3) from 3.0 x 10 -4 to 7.5 × 10 -4. On the other hand, M value for 2.25 Cr-1 Mo steel as reported in Fig. 7 is 6 x 10 -4 which is lower than the expected value of 7.5 x 10 -4. A difference in the ~ and the/'values of the two steels, could account for the difference observed between the expected and the calculated M values.

Creep l(fe prediction

The discussion in the previous three sections shows that the experimental values: K = 0.8 in equation (4) and that of the M equal to 3 x 10 -4 in equation (3) for the 1 Crq).3 Mo-0.25 V steel have some physical bases. Therefore, the proposed equation (3) could be used to make a first order estimate of the creep life of components which experience stresses around 50 MPa and temperatures around 550°C.

The following steps can be suggested for a rapid assessment of the remaining creep life of a component made of 1 Crq3.3 Mo 0.25 V steel and subjected to the usual service temperatures < 550~C (barring acci- dental short term excursion to a higher temperature).

1. Determine the hot yield strength, aT, at the service temperature and estimate OH from the relationship

a H : 0.48 Oy.

2. Determine ~m at six stress levels, three of them between 0.4 to 0.5 times and the other three at 0.6, 0.7 and 0.8 times the hot yield stress. The maximum duration of one such test is 1500 h only.

3. Plot (gm) °2~ VS a curve in a manner similar to that in Fig. 4 and determine a n.

4. If OH is greater than the service stress, a, calculate the residual life in terms of the time required to produce the balance amount of strain to make up the maximum permissible strain of 1%. The calculation uses the simple relationship

(~m) 025 : 3 × 10 -4 a .

5. If a n is less than the service stress, a, retire and replace the component at the earliest.

Five significant points should be noted with regard to the above procedure. First, more creep data, particularly at higher temperatures, lower stresses and for the other ferritic steels should be generated to validate the procedure further. Second, the total duration of creep testing required for life prediction which is about 15,000 h at present, can decrease five fold if the procedure proposed above is used. Third, in future, after the equations (2) and (3) are verified by data obtained from material exposed to service exposure longer than 105 h (say 20 or 30 years), the creep testing can be reduced still further or even entirely dispensed with. At this point, the actions proposed at steps 2 and 3 can be avoided. Fourth, the strain experienced by the component in service should be monitored using an appropriate procedure.

4. CONCLUSIONS

1. An analysis of the creep data generated at 550°C, on the 1 Cr-0.3 Mo-0.25 V steel in the five different microstructural states, shows that the steel exhibits a threshold stress above which "runaway" creep could take place. This threshold stress depends on the microstructure.

2. The threshold stress value decreases by only a small amount as a result of 105h service exposure at 540:C/50 MPa.

3. At applied stress values less than the threshold value, the minimum creep rate has an unique relationship with the applied stress, for all the five microstructural states of the steel investigated.

4. The threshold stress can be more readily estimated from the hot yield stress since the yield stress is about twice the threshold stress.

5. The concept of the threshold stress can be used for a rapid assessment of the remaining life, using a simple analysis of the steady state creep data.

REFERENCES

1. J. D. Parker and B. Wilshire, Metals Sci. 9, 248 (1975). 2. P. W. Davies, G. Nelmes, K. R. Williams and B.

Wilshire, Metals Sci. 7, 87 (1973). 3. D. A. Miller, Mater. Sci. Engng 54, 169 (1982). 4. D. A. Miller, W. J. Plumbridge and R. A. Bartlett.

Metals Sci. 15, 413 (198l). 5. B. Reppich, H. Bugler, R. Leistner and M. Schutze.

Creep Fracture of Engineering Materials Structures. Proc. Int. Conj. Univ. College, Swansea, U.K. (edited by B. Wilshire and D. R. J. Owens), p. 279. Piperidge Press (1984).

6. R. A. Stevens and P. E. O. Flewitt, Acta metall. 29, 867 (1981).

7. H. E. Evans and G. Knowles, Acta metall. 25, 693 (1977).

8. S. Purushothman and J. K. Tien, Acta metall. 26, 519 (1978).

9. T. Lagneborg and B. Bergman, Metals. Sci. 10, 20 (1976).

10. J. D. Parker and B. Wilshire, Metals. Sct. 9, 248 (1975).


I1. W. J. Evans and G. F. Harison, Metals. Sci. 10, 317 (1976).

12. K. R. William and B. Wilshire, Mater. Sci. Engng 28, 289 (1977).

13. R. Singh and S. Banerjee, Scripta metall, mater. 24, 1093 (1990).

14. R. Singh and S. Banerjee, Mater. Sci. Engng A 132, 203 (1991).

15. H. D. Wolf, thesis, Univ, of Erlangen, Nurnberg, Germany (1990).

16. W. Blum and B. Reppich, Creep Behaviour of Crys- talline Solids, Vol. 3 of Series Prog. in Creep and Fracture (edited by B. Wilshire and R. W. Evans), p. 83. Univ. College, Swansea, U.K. (1985).

17. G. Sundarajan, Trans. I IM 42, S 189 (1989). 18. B. F. Dyson and M. Mclean, Acta metall. 31,295 (1983). 19. S. G. R. Brown, R. W. Evans and B. Wilshire, Mater.

Sci. Engng 84, 147 (1986). 20. M. F. Ashby and B. F. Dyson, Advances in Fracture

Research (edited by S. R. Valluri, et al.), Vol. 1, p. 30. Pergamon Press, Oxford (1984).

21. M. McLean, Acta metall. 33, 545 (1985). 22. R. C. Ecob and H. E. Evans, Acta metall. 35, 805 (1987). 23. V. Foldina, A. Jakobova, T. Prinka and J. Sabotk,

Proc. Conf. on Creep Strength in Steel and High Tem- perature Alloys, Sheffield, 1972, pp. 230-236. The Metall. Soc., London (1973).

24. B. J. Cane and R. D. Townsend, Flow and Fracture at Elevated Temperatures (edited by R. Raj), pp. 279 316. Am. Soc. Metals, Metals Park, Ohio (1983).

25. H. Weber, Report No. 35-V, Mannesman, West Germany.

26. R. R. Jensen, D. thesis, Columbia Univ. (1984). 27. J. C. Gibbeling and W. D. Nix, Metals Sci. 11, 435

(1977). 28. M. McLean, Proc. R. Soc., Lond. A371, 279 (1980).

29. H. Burt, J. P. Dennison and B. Wilshire, Mater. Sci. J. 13, 295 (1979).

30. J. D. Whitternberger, Metall. Trans. 15A, 1753 (1984). 31. K. Kucharova, A. Orlova, A. Otkawa and K. Cadeel,

Mater. Sci. Engng, in press. 32. J. Peterseim and G. Santhoff, Steel Res. 57, 19 (1986). 33. E. Artz and M. F. Ashby, Scripta metall. 16, 1285

(1982). 34. W. J. Evans and G. F. Harrison, Metals Sei. 10, 307

(1976). 35. M. Mclean, Proc. R. Soc. Lond. A373, 93 (1980). 36. B. Bergman, Scand. J. metall. 4, 97 (1975). 37. H. Burt, J. P. Dennison and B. Wilshire, Mater. Sci. J.

13, 295 (1979). 38. A. Artz, Creep o f Dispersion Strengthened Materials: A

Critical Assessment, Res. Mechanica, in press (1990). 39. R. Russel, R. R. Jenson and J. K. Tien, Metall. Trans.

16A, 1049 (1985). 40. K. R. William and B. J. Cane, Mater. Sci. Engng 38, 199

(1979). 4l. V. C. Nardone and J. K. Tien, Scripta metall. 20, 797

(1986). 42. L. M. Brown, Proc. 3rd Ris. Symp. on Metall. Mater,

Sci. Ris, Denmark (1982). 43. J. Rosier and E. Artz, Acta metall. 36, 1043 (1988). 44. J. Rosier and E. Artz, Acta metall. 38, 671 (1990). 45. A. M. Brown and M. F. Ashby, Scripta metall. 14, 1297

(1980). 46. M. A. Krishtal, Diffusion Process in Iron Alloys, p. 135.

Israel Prog. for Sci. Translation, Jerusalem (1970). 47. J. H. Brophy, R. M. M. Rose and J. Wulff, The

Structure and Prop. of Metals, Vol. 2: Thermodynamics o f Structure p. 194. Wiley Eastern Pvt. Ltd, New Delhi (1966).

48. D. Hull, Introduction to Dislocation, p. 222. Pergamon Press, Oxford (1965).

Resisting stress of a low alloy ferritic steel after creep exposure in service

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