Evolution of microstructure in a modified 9Cr–1Mo steel during short term creep

Materials Science and Engineering A245 (1998) 285 – 292

Evolution of microstructure in a modified 9Cr–1Mo steel during short term creep E. Cerri a,*, E. Evangelista a, S. Spigarelli a, P. Bianchi b a

INFM/Dipartimento di Meccanica, Uni6ersita` di Ancona, Via Brecce Bianche, 60131 -Ancona, Italy b ENEL-CRAM, Via Volta 2, Cologno, Milano, Italy Received 15 May 1997; received in revised form 19 September 1997

Abstract An investigation of the effect of creep exposure on the microstructure of a 9Cr – 1Mo alloy for steam tubing was performed. The samples were machined from a tube, austenised at 1323 K for 15 min and air cooled to room temperature, followed by tempering at 1023 K for 1 h. Creep tests were performed at 848, 873, 898 and 923 K for different loading conditions. The conventional power law was used to describe the minimum creep rate dependence on applied stress; the stress exponent was found to increase when temperature decreased. Transmission electron microscopy (TEM) of the crept samples showed that during creep both subgrain and particle size increased; the statistical analysis of the dimensions of the precipitates revealed a bimodal distribution of particles that coarsen during creep exposure at testing temperatures. A linear dependence of subgrain size on the inverse of the modulus compensated stress was used to describe the softening of the dislocation substructure. A similar relationship was found to be also valid for particle carbides. © 1998 Elsevier Science S.A. All rights reserved. Keywords: 9Cr – 1Mo steel; Creep; Microstructure

1. Introduction The T/P91 steel (a 9Cr – 1Mo modified by Nb and V addition) is a relatively new structural alloy that was originally developed for use as a steam generator material for advanced fast breeder reactors and also used widely in the power generation industry for tubing applications for prolonged service at temperatures of : 873 K. All these uses require knowledge of the influence of long-term high temperature exposure on the tempered martensitic microstructure and on mechanical properties [1]. It has been reported that the creep properties of such steels are influenced by the dispersion of the carbide and carbo-nitride phases, consisting mainly of M23C6, Nb(C,N) and VN precipitates [2,3]. These precipitates produce a hardening effect by preventing the movement of subgrain boundaries, by impeding knitting reactions between free dislocations and subgrain boundaries and by pinning subboundary dislocations [4 – 6]. Creep exposure produces the recovery of dislocations, the ag* Corresponding author. E-mail: [email protected] 0921-5093/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S0921-5093(97)00717-X

glomeration of carbides and the growth of martensite laths subgrains [7–9]; while all these phenomena result in a gradual loss of creep strength, the additional precipitation of fine MX (M stands for a metal and X for C or N), that is frequently observed during creep, could reduce the microstructural instability by obstructing the dislocation mobility. In order to clarify the effect of substructure evolution on creep behaviour, there has been an effort to separate the subgrain contribution to the strengthening of the material from the particle contribution in various oxide dispersion strengthened materials [10]. In martensitic 9Cr–1Mo steel Straub et al. [11] succeeded in a microstructurally based simulation of the long-term creep behaviour including all the microstructural parameters. The aim of the present study is to determine the microstructural evolution of a T91 steel. A knowledge of the strengthening and degradation mechanisms is an important prerequisite for the development of high creep resistant chromium steels. In particular, subgrain and particle size evolution will be followed by transmission electron microscopy (TEM) after short term creep exposure tests at 848, 873 and 923 K.

E. Cerri et al. / Materials Science and Engineering A245 (1998) 285–292

286 Table 1 Chemical composition of the steel C 0.091

Mn 0.46

Si 0.40

Cr 8.76

Mo 0.94

N 0.064

Ni 0.11

Nb 0.07

Al 0.35

V 0.19

Values are in wt.%.

Fig. 1. Modified 9Cr –1Mo steel after normalising and tempering: (a) optical micrograph illustrating the microstructure; (b) TEM micrograph; (c) size distribution of the precipitates (the experimental data in form of histogram and cumulative frequency, ’, are compared with the calculated distribution, ——).

2. Material and experimental techniques The material investigated in this study was provided by Dalmine (Dalmine, Bergamo, Italy); the chemical composition is given in Table 1. The heat treatment consisted of austenitising at 1323 K for 15 min followed by air cooling and tempering at 1023 K for 1 h. The hardness after thermal treatment was 20.2 HRC. Creep specimens 5 mm in diameter and 50 mm in gauge length were machined directly from a tube. Short term constant load creep tests were performed at temperatures of 848, 873, 898 and 923 K, under stresses ranging

from 85 to 240 MPa. Specimens were usually tested until failure, with the only exception of a test carried out at 873 K–130 MPa and interrupted after 504 h (roughly 1/5 of the creep life). The elongation of the sample was measured by LVDT (linear variable displacement transducers) and continuously recorded. The metallographic investigation performed in this work included hardness measurements, optical microscopy (OM) and TEM. In particular, TEM was used for studying the particle size distribution in the matrix by using an image analyser on TEM negatives. The magnifications used were 22000, 45000 and 75000×

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for particles. The lower limit of particle size detection is 10 nm. For observations, slices 1 mm thick were cut from the specimens at 45° to the longitudinal axis and mechanically ground to 100 – 150 mm. Disks were then cut from the slices and electrochemically polished in a solution of 10% HClO4 in methanol at 238 K and 20 V. The observations were carried out in TEM Philips CM12 at 120 kV. OM samples were prepared by hot mounting in acrylic moulding powder followed by grinding to 120 grit and polishing to 0.25 mm cloth. Specimens were etched using Picral reagent. Subgrain size was measured by the intercept linear method.

Fig. 3. Hardness after creep as measured on samples tested at 873 K; the moderate hardening at 200 MPa could reflect the presence of a dispersion of fine particles precipitated during creep. The softening observed for lower stresses corresponds to the coarsening of the precipitates and recovery of the dislocation substructure.

3. Results

3.1. Initial microstructure Fig. 1(a) shows the optical micrograph of the heat treated sample, before creep, as taken from the tube. The structure is the one typical of tempered martensite, consisting of parallel martensite laths arranged inside a prior austenite grain. A TEM micrograph of the same sample illustrates (Fig. 1(b)) the presence of martensitic laths containing a high density of free dislocations. The substructure is partially organised in subgrains and the particles precipitated during tempering are mainly placed on tempered martensite lath boundaries and prior austenite grain boundaries. These particles are M23C6 carbides and have a mean equivalent diameter of :80 nm after tempering. Their gaussian distribution is illustrated in Fig. 1(c). The phase is based on Cr23C6 type, but it can dissolve a multitude of other elements (Fe, Mo) [12,13]. The subgrain size of the as-tempered steel has been determined to be 0.41 mm.

3.2. Creep results

Fig. 2. Mechanical results on 9Cr–1Mo steel. (a) Strain rate o; – time t curve recorded at 873 K. (b) Minimum strain rate (o; m) dependence on applied stress s.

The strain rate o; as a function of time t under different stresses at 873 K is reported in Fig. 2(a). The secondary region is rather short over the investigated range of stress and temperature. Fig. 2(b) illustrates the minimum creep rate (o; m) as a function of the applied stress s. The o; m dependence on s can be conveniently described by the following equation:

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Fig. 4. TEM microstructure after creep at different loads and temperatures: (a) 873 K – 175 MPa, (b) 873 K – 130 MPa, (c) 923 K – 160 MPa and (d) 923 K – 115 MPa.

o; m = A(T) s n,

(1)

where A(T) is a temperature dependent parameter and n is the stress exponent. The stress exponent was found to decrease from 8.7 to 6.0 as temperature increases from 848 to 923 K. Further, the separation of the lines in Fig. 2(b) clearly indicates that Q (Q = [( ln o; m/ ( ln(− 1/RT)] at constant s) increases when stress decreases. Typical values of the apparent activation energy are Q = 726 and 600 kJ mol − 1 for s= 180 and 220 MPa, respectively. Such values are significantly higher than the activation energy for self diffusion (250 kJ mol − 1) [14]. The hardness in creep tested specimens as a function of the stress is reported in Fig. 3 (873 K). Creep exposure produces a pronounced microstructural softening, while the fracture elongation remains essentially constant in the investigated range of stress.

3.3. Microstructure after creep TEM micrographs illustrating the microstructural evolution during creep at 873 K are presented in Fig. 4(a) and (b) for different values of the stress. It is

shown that the subgrains grow during creep with respect to the initial state (Fig. 1) and there is a further rearrangement of free dislocations. The particles, depending on their dimensions, in general try to pin the subboundaries, but when the subgrains grow, some subgrain boundaries disappear leaving some of the particles in the subgrain interior as shown in Fig. 4(b). Consequently the fraction of particles attached to subgrain boundaries decreases during creep with respect to the initial one [15]. At 923 K, the subgrain size is larger at any experimental load considered than at the lower temperature (Fig. 4(c) and (d)). The particle distributions in creep tested specimens are illustrated in Figs. 5 and 6. At 848 and 873 K, and even at 923 K at low stresses, the particle population is bimodal, containing small particles that were not present in the as-received sample and that can be identified to be of MX type, and particles of larger dimensions already present in the initial state (M23C6 type). An attempt was made to separate the two types of precipitates using the same procedure suggested by Straub et al. [11]. The distribution curve is considered to be the sum of two contributions, corresponding to large (I) and fine (II) particles; then the cumulative

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Fig. 5. Size distribution of the precipitates observed after creep at 873 K (the experimental data in form of histogram and cumulative frequency, ’, are compared with the calculated distribution, ——): (a) 873 K – 200 MPa, (b) 873 K – 175 MPa, (c) 873 K – 130 MPa and (d) after 504 h at 873 K– 130 MPa.

distribution F can be expressed as F = FIF I + FIIF II, where F I and F II are the single cumulative distributions, while FI and FII are the numerical fractions of the two families of precipitates (FI +FII =1). This procedure permits the approximation of the gaussian curve for each type of precipitate, and to calculate a cumulative distribution that, in general, is in good agreement with the experimental one. The effect of decreasing stresses (the creep time to rupture is longer) on particles is to promote the growth of both types of particles, as can be seen in Fig. 5, even if with different kinetics. Fig. 6, on the other hand, confirms that precipitation of fine particles occurs at 923 K only for low stresses (i.e. long time of exposure).

4. Discussion It is now widely accepted that the microstructural evolution of modified 9Cr steel is extremely important since it affects the creep strength of the material; in the base 9Cr–1Mo steel, for example, tempering produces a dispersion of M23C6 precipitates that stabilise the

microstructure by pinning subgrain boundaries and so reduce the effectiveness of recovery mechanisms [16– 19]. Coarsening of M23C6, then, represents one of the most important degradation processes in these steels. In Nb and V containing 9Cr–1Mo alloy, fine and dispersed MX phases precipitate during tempering and/ or creep, and represent an additional source of strengthening. Since the presence of these fine Nb and V carbo-nitrides and nitrides, as well as the coarsening of M23C6, strongly affect the creep strength [20], it seems useful to analyse the microstructural evolution in the steel.

4.1. Precipitation of new phases Foldyna et al. [16,17,21–23] made an extensive effort in order to clarify the effect of chemical composition and heat treatment conditions on the precipitation of new phases (namely VN, Nb(C,N) and AlN) in modified 9Cr steels. They showed, for example, that in order to optimise the creep strength, one of the most important parameters to be considered is the amount of N in solid solution (NSS) after austenitising and temper-

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ing [24]. A high amount of N in solid solution and a sufficient content of nitriding forming elements (mainly V), in fact, permit the precipitation of fine MX that obstruct dislocation mobility; a high Al content, from this point of view, is detrimental, since AlN precipitates in the form of coarse particles and reduces the N available for precipitation of fine VN [23]. Similarly, large undissolved MX do not give any contribution to the creep strength; in Fig. 7(a), the solubility curve for VN at 1323 K, as reported by Foldyna et al. [24], is compared with the chemical composition of the present

steel (the effect of Al and Nb content is neglected here). Examination of the figure reveals that the normalising treatment should maintain V and N in solid solution. Yet, the real situation appears to be somewhat more complicated, since the presence of Al and Nb can reduce, as previously observed, the N available for precipitation [22,24]. Fig. 7(b) shows the effect of chemical composition (Nb, V and Al content) on the amount of N in solid solution after normalising [24]; for a steel containing 0.19% V and 0.07% Nb, then, NSS after normalising should be close to 0.06% (:86% of the total content in the model steel). Subsequent tempering could produce the precipitation of fine MX; such extensive precipitation was observed by Straub et al [11] in a P91 steel after tempering at 1033 K (12 h)+ 1023 K (8 h), while (Fig. 1(c)) only a very low fraction of extremely fine particles of dimension close to 10 nm was detected in our steel in the as-received condition. The solubility curve in ferrite at 1023 K, given by Foldyna et al. [24], indeed, suggests that a relatively large amount of MX should precipitate during tempering at 1023 K; on the other hand, a tempering treatment of 1023 K (1 h) is probably too short to produce the extensive precipitation observed by Straub et al. [11]. The kinetics of precipitation is undoubtedly faster during creep, due to the lower temperatures (i.e. to the lower solubility products), and to the presence of a high dislocation density that provides a large number of nucleation sites. Comparison of Fig. 5(a) and (d) suggests that the microstructure after 504 h at 873 K–130 MPa is equivalent to the one observed in a sample tested for 35.6 h at 873 K–200 MPa; then it can be concluded that not only nucleation, but also growth and coarsening of the different types of precipitates are strongly influenced by the applied stress.

4.2. Degradation processes In order to characterise the evolution of subgrain size during creep exposure, the dependence of subgrain size l as a function of the inverse of modulus compensated stress (s/E) is illustrated in Fig. 8(a). The general relation is: l: K%b(s/E) − p,

Fig. 6. Size distribution of the precipitates observed after creep rupture at: (a) 923 K–115 MPa, (b) 923 K–85 MPa, and (c) 848 K– 185 MPa.

(2)

where b is the Burgers vector= 0.247 nm [20], K% is a material constant, l (mm) is the subgrain or cell size at the rupture, E is the Young’s modulus and p is a constant that depends weakly on the material, generally close to 1 [25]. In a-Fe, this relationship is substantially fulfilled, even if an increase in the exponent p is observed for stresses larger than 200 MPa at 873– 973 K. Further, a weak dependence of p on temperature is observed [26].

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Fig. 7. (a) Solubility curve for VN at 1323 K for a 9Cr–1Mo steel (’ shows the studied steel). (b) Effect of composition of two 9Cr –1Mo model steels on the amount of N; NSS element in solid solution; NNb element combined with Nb; NAl element combined with Al (see [24]).

In the present study, any dependence on temperature was neglected; interpolation of the experimental values of subgrain dimension gives K% =1.39 and p= 1.09. The agreement between experimental data and calculated curve is tolerable, and then it can be accepted that the subgrain size (at rupture) is a function of the flow stress divided by modulus experienced by the sample during creep at a fixed temperature. The linear relationship, Eq. (2), has been successfully used by Sherby et al. [10] to relate the strain rate to microstructural parameters like the subgrain and particle sizes and their distributions (the overall relationship refers to constant structure creep tests). An empirical expression of similar form can be used to correlate the particle size with stress, i.e.: dM23C6 :K¦b(s/E) − b,

microstructural evolution, and of the effect on the creep response, will be given elsewhere (unpublished research). Even if it is widely accepted that the subgrain size is determined by applied stress, a major question is to determine if the presence of particles has a direct influence on it. It is reported that there is no indication for a direct influence of particles on the subgrain size [27], but Eqs. (2) and (3) show that it is possible to correlate the subgrain size with the dimension of the precipitates (Fig. 8(b)). Such a relationship suggests that the coarsening of the subboundary particles promotes the subgrain growth; this result, in turn, supports the idea that the main effect of M23C6 is to pin subgrain boundaries, reducing the effectiveness of the recovery mechanisms.

(3)

where K¦ =3.6 and b = 0.68 nm. This expression permits the extrapolation of the dimensions of particles in samples crept under different loading conditions and so gives a rough idea of their distributions. Eq. (3) is probably the result of the dependence of test time on stress; then, since the dimension of the particles is a function of time of exposure [11,16,17] one can obtain an indirect dependence of particle size on stress. An accurate estimation of the particle dimension at a given time and under specific stress and temperature requires a detailed analysis of the influence of all these parameters on the kinetics of the processes. The different growth and coarsening rates, for example, are connected with the dislocation density and of course with the diffusion coefficient of solute atoms in the matrix surroundings. A detailed study of these aspects of the

5. Conclusions An evaluation of the microstructural evolution of a T91 crept steel has been performed under different loading and temperature conditions by TEM. The following conclusions were observed. “ The particle distribution is bimodal in most of the creep tested specimens, including a population of smaller dimensions (MX carbides and nitrides) precipitated during testing and also M23C6 carbides mainly located on subgrain boundaries. Both the types of precipitates coarsen during creep. “ The subgrain size depends linearly on the inverse of the modulus compensated stress for tests to failure; also the particles exhibit a dependence on (s/E), leading to the possibility to correlate the subgrain

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for funding the research, respectively. A further acknowledgement is due to Professor M.E. Kassner for useful discussion.

References

Fig. 8. a) Subgrain and M23C6 size as a function of applied stress; the error bars indicate the S.D. of the measured values (subgrain size), and the S.D. as obtained from the distribution curves in Figs. 5 and 6 (particle size). (b) Relationship between M23C6 and subgrain size at failure.

size and the M23C6 dimensions. This, in turn, could reflect the effect of degradation processes (coarsening of the subboundary precipitates) on the softening of the microstructure.

Acknowledgements The authors would like to thank Dalmine (Bergamo) and Enel-Cram (Milano) for providing the material and

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