Microstructural evolution in a PH13-8 stainless steel after ageing

Acta Materialia 51 (2003) 101–116 www.actamat-journals.com

Microstructural evolution in a PH13-8 stainless steel after ageing Z. Guo a, W. Sha a,∗, D. Vaumousse b a

School of Civil Engineering, Queen’s University of Belfast, Belfast, UK b Department of Materials, University of Oxford, Oxford, UK

Received 6 May 2002; received in revised form 2 August 2002; accepted 2 August 2002

Abstract The precipitation process in a wrought PH13-8 steel during ageing was investigated using a position sensitive atom probe. The precipitates formed are enriched in Ni and Al, and depleted of Fe and Cr, yet the composition is far from the stoichiometric NiAl phase. They may take on different shapes at different temperatures. The hardening effects observed during early stages of ageing should be due to the redistribution of atoms such as Fe, Cr, Ni, and/or Al. Particle coarsening takes place simultaneously with the development of the composition of the NiAl-enriched precipitates. Mo segregation at the precipitate/matrix interface was not found.  2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. Keywords: Atom probe; Microstructure; Ageing; Precipitation hardening; Stainless steels

1. Introduction PH13-8 stainless steel is a martensitic precipitation hardening (PH) steel. It has high strength and hardness with good levels of resistance to both general corrosion and stress-corrosion cracking. In addition, the alloy exhibits good ductility and toughness in large sections in both the longitudinal and transverse directions, and offers a high level of useful mechanical properties under severe environmental conditions superior to PH17-4 and PH15-5 stainless steels. It has been used for many

Corresponding author. Tel.: +44 28 9027 4017; fax: +44 28 9066 3754. E-mail address: [email protected] (W. Sha). ∗

applications, such as landing gear parts, nuclear reactor components and petrochemical applications requiring resistance to stress-corrosion cracking [1]. A number of studies have shown that the ferritic and martensitic phases in high alloy steels (including stainless steels such as PH13-8 steel) containing nickel and aluminium can be hardened by ageing at temperatures above 400 °C [2–5]. The strengthening is reported to be due to the precipitation of the ordered phase NiAl which has a B2 (CsCl) superlattice structure. However, due to the instrument limitation, in the past, it was not possible to analyse directly fine precipitates of nanometer scale. Therefore, the materials studied had to be treated to seriously-overaged conditions so as to achieve precipitates of detectable sizes by

1359-6454/03/$22.00  2002 Acta Materialia Inc. Published by Elsevier Science Ltd. All rights reserved. doi:10.1016/S1359-6454(02)00353-1

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transmission electron microscope (TEM) [6,7]. The most recent study with wrought PH13-8 stainless steel was carried out by Seetharaman et al. in early 1980s [6]. They found that precipitates formed during ageing at temperatures lower than 500 °C for 4 h could not be resolved with TEM. When samples overaged at 575 °C were studied, the precipitates were claimed to be spherical, and uniformly distributed in the matrix. The selected area diffraction pattern (SADP) analysis revealed the existence of the intermetallic compound NiAl of B2 structure. Another relevant study was carried out by Taillard et al. on Fe-19wt%Cr-Ni-Al systems [8,9]. They found that the NiAl precipitates were homogeneously nucleated and coherent with the matrix after 400 min at 650 °C. It should be noted that the precipitation information obtained in seriouslyoveraged conditions does not necessarily represent what happens in the commercially treated materials, where the ageing treatment is usually 1 or 4 h at temperatures below 600 °C. The most recent work on PH13-8 stainless steel is on a cast grade [7]. The precipitate type was reported as NiAl through SADP analysis and observed to be spherical (H1150M treatment, see Section 2.1 for details). It was very difficult to image the precipitates using TEM even after ageing at 510 °C for 4 h. The size only reaches 40–50 nm after the H1150M treatment. It is worthwhile mentioning that the age hardening kinetics of the cast grade differs from its wrought counterpart significantly [10]. Other previous work was also carried out on overaged materials, as otherwise the precipitates were not large enough to allow TEM examination and SADP identification [11,12]. How and when the precipitates start to form during ageing of the PH13-8 alloy, and how they evolve in terms of size, composition and even type remain unclear. Atom probe field ion microscopy (APFIM) proved to be very powerful in the investigation of small precipitates when other microscopy techniques are unsatisfactory [13,14]. Its unique capability of measuring composition variations in a nanometer scale, together with equal detection efficiency for all elements, make it particularly suitable for investigation of early stages of precipitation. The early atom probe (AP) analysis was essentially one dimensional in nature. Recent

developments allow a 3-dimensional reconstruction of the distribution of different atoms in the analysed volume [15]. Atom probe (AP) microanalysis was carried out in determining the composition of the precipitate in an Fe-20Cr-2Ni-2Al (at%) alloy [16]. It was found that the composition of the B2 intermetallic phase is close to the stoichiometric composition NiAl with only a limited amount of dissolved iron, which decreases with ageing time. The treatments used in their study were ageing at 550 °C for 6, 17 and 117 h. Miller and Hetherington studied an Fe-Ni-Al-Mo system [17]. They found that Mo partitioned preferentially to the α matrix, and NiAl β’ precipitates also contained approximately 11% Fe. Some previous work with Cr-containing steels is summarised here since the studied PH13-8 steel contains chromium. Most of the alloys previously studied contain a higher amount of Cr than PH138. Alloys Fe-30.1Cr-9.9Co, Fe-20.2Cr-8.8Al0.55Ti, and Fe-26Cr-(0, 3, 5, or 8)Ni (all in at%) all show spinodal decomposition of Cr during the ageing process [18–20]. Stainless steel PH17-4 (0.3C-17.5Cr-4Ni-3Cu-lSi in at%) was found to be strengthened by ⑀-Cu particles produced during ageing [21]. Murayama et al. [22] found that the phase separation occurs through a spinodal mechanism while ageing at 400 °C. Tempering this alloy at 580 °C for 4 h does not lead to any Cr separation in the martensite phase, which is the condition before this alloy enters service. Another maraging stainless grade 1RK91 (13Cr-9Ni-2Mo-2Cu in wt%) is of low-Cr type. No α-α’ transformation, i.e. the spinodal decomposition of Cr, was found in this alloy at 475 °C even after 1000 h [23]. The good mechanical properties are attributed to the heterogeneous precipitation of Ni-rich/(Al-Ti) precipitates, possibly Ni3(Ti,Al) type, on the Cu particles formed in the very early stages of ageing [24]. Enrichment of Mo at the precipitate/matrix interface impedes the coarsening of particles during ageing. An aged ferritic-martensitic steel (with 11 at% Cr) was investigated using atom probe [25]. After ageing at 400 °C for 17000 h, long range fluctuations with composition difference of 2.9 at% were present in the alloy. From the above summary of the previous work, it is clear that there is lack of information in a few

Z. Guo et al. / Acta Materialia 51 (2003) 101–116

aspects about PH13-8 steel during ageing, such as, when and how the precipitates start to form during ageing; how they evolve in terms of type, composition, shape, size and fraction, and what kind of precipitates contribute to the strengthening effects. It will also be useful to determine whether Cr decomposes spinodally in a steel of such a low Cr level. If so, how fast does this take place? What will be the influence on the evolution of other precipitates and strengthening kinetics? Does Mo segregate to the precipitate/matrix interface, as reported for the 1RK91 steel? Answers to the above questions will help in understanding the strengthening mechanisms during ageing, and make possible a computer model of the age hardening kinetics based on existing precipitate hardening theories [26,27].

2. Experimental procedures

2.1. Material

103

Table 2 Ageing treatments of the PH13-8 stainless steel Tage 510 °C tage ∗ #

593 °C

4 min 15 min 40 min 4 h#

6 min 30 min∗ 1 h

Peak hardness. Age hardening curves are given in Fig. 1. The commercial H950 treatment.

650 and 720 °C. Blocks of size about 12 × 12 × 12 mm3 were cut from the as-received material. They were then reheated at 927 °C for 1.5 h followed by water quenching, with no refrigeration. The solution-treated blocks were put into a furnace with temperature 510 °C or 593 °C for ageing treatment. They were aged at 510 °C for 4, 15, 40 min and 4 h, or at 593 °C for 6, 30 min and 1 h, Table 2, where Tage denotes ageing temperature and tage, ageing time. A layer of 2 mm was removed from the sample surface to avoid the possible influence of oxygen when specimens for atom probe were prepared.

The typical ageing treatment for PH13-8 steel is ageing the solution-treated material at intermediate temperatures, for instance, 510 °C for 4 h (H950), 593 °C for 4 h (H1100), or 760 °C for 2 h followed by 4 h at 621 °C (H1150M). Such treatments cause hardening by precipitation of intermetallic compounds and provide a good combination of strength and ductility [1]. The material studied was provided by Allvac Ltd., UK. Its composition in both wt% and at% is given in Table 1. The material was vacuum induction melted plus vacuum arc remelted (VIM/VAR) to a 500 mm ingot, homogenised at 1250 °C for 12 h, forged from 1100 °C to 280 mm diameter, reheated to 1150 °C and forged to 146 mm diameter, cooled to below 90 °C, and annealed between

Fig. 1.

Age hardening of PH13-8 steel.

Table 1 Chemical composition of the PH13-8 Mo steel

wt% at%

C

Al

Cr

Mo

Ni

Ti

Si

Co

Mn

Zr

P

S

N

Fe

0.03 0.19

0.97 2.02

12.43 13.36

2.15 1.36

8.39 7.96

0.067 0.184

0.07 0.17

0.01 0.01

0.02 0.02

0.004 0.002

0.006 0.005

0.002 0.004

0.001 0.004

bal. bal.

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2.2. Sample preparation methods For examination in the atom probe, the material is required to be in the form of sharp needles of ⬍ 100 nm end radius. The form is achieved by cutting blanks on a diamond saw, then by a standard two-stage electrolytic polishing. The specimen is mounted in a nickel tube before introduction into the atom probe. For a description of the operation and principles of the atom probe field ion microscopy, the reader is referred to the book by Miller and Smith [13]. As aforementioned, the limitations of the early atom probe analysis were largely overcome by efforts originated by Cerezo et al., who re-designed the atom probe to incorporate a large area, based on a position sensitive ion detector system. The instrument, termed a Position Sensitive Atom Probe (PoSAP), has found its applications in many areas of materials science since then [28], and is used in the current study. The specimens were first imaged in field ion microscope (FIM) mode in neon and then examined at a temperature of 65 K in the atom probe. The first step (FIM) is used to monitor the development of a stable, approximately smooth, hemispherical tip for subsequent analysis. Where there is image contrast, FIM can also be used to select the analysis positions that might be of particular interest. This was not the case in the current study, due to the lack of image contrast between the precipitates and the matrix. 2.3. Interpretation of atom probe data The most straightforward way is to use the atom map, the elemental atomic distribution of the studied material reconstructed from atom-by-atom raw data. In the atom map display, the size of the atoms can be set arbitrarily, usually smaller than the real atomic diameter. In this paper, some variation of the atom size has been used to achieve the best visual effect. However, this display cannot reveal the slight heterogeneous atom distribution with no obvious particles. The experimental atom probe data can be statistically analysed to establish physical parameters. Such methods have been reviewed by Hetherington and Miller [29], and may also be referred to in Refs. [13] and [14]. In order to

extract useful information, it is convenient to divide the data into samples, each containing an equal number of ions, for which the composition is calculated. The sample size or block should be about 50–200 [13]. In the current study, blocks of 100 atoms have been chosen. All concentration values are quoted in atomic percent, unless otherwise stated. The errors on the concentration values are estimated by the standard deviation, √c(1 ⫺ c) / N, where c is the measured concentration and N the total number of detected atoms [13].

3. Result analysis The hardness measurement was carried out on a Mitutoyo HM-124 machine, with 2 kg working load. The age hardening curves at 510 and 593 °C of the studied material are plotted in Fig. 1. The hardness reaches peak after 30 min at 593 °C, but still increases after ageing for 4 h at 510 °C. The ageing behaviour is similar to what was obtained before for a wrought PH13-8 stainless steel [1], but significantly different from that of a cast PH13-8 steel [7,10]. The PoSAP machine was first operated in FIM mode to allow the development of a good tip for probe analysis. In field ion images of maraging and PH steels, austenite normally shows a contrast. However, no contrast was observed in the experiments, Fig. 2. Atom probe data is obtained from the instrument in the form of a sequence of the mass-to-charge ratio of each of the ions evaporated. The PoSAP instrument used allows a much higher mass resolution than its precursor, through the introduction of a reflectron-based energy-compensation system. The mass spectrum of the specimen aged for 4 min. at 510 °C is shown in Fig. 3 as an example. A minor problem here is the exact coincidence of 58 Fe2+ and 58Ni2+, and 64Ni2+ and 96Mo3+ peaks in the mass spectrum. In the present study, these are regarded as 58Ni2+ and 96Mo3+, respectively, after consulting the abundance distribution of each isotropic species in natural form and the amounts of the elements in the alloy. A small peak at Mass/charge ratio 1 was also found, corresponding to H+. This is not included in the composition calculations, as the hydrogen atoms are believed to

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the as-received alloy. Frequency distribution (FD) analysis was carried out to examine whether each element distributes randomly. If all the atoms of one type of element are of random distribution, the frequency distribution will follow what is described by a binomial model. FD analysis result deviating significantly from the binomial model means that the element distributes non-randomly. FD analysis reveals that Cr, Fe and Ni are not randomly distributed in the matrix whereas the other elements are. At present, it is difficult to identify which of the following is the cause for this inhomogeneity:

Fig. 2. FIM image of a PH13-8 steel sample aged at 593 °C for 30 min. The distance across the image is about 70 nm.

Fig. 3. Mass spectrum of the sample aged for 4 min at 510 °C.

앫 There was inhomogeneity in the as-received material that was not erased after the reaustenisation treatment (1.5 h at 927 °C); 앫 The sample was homogeneous at the end of the re-austenisation treatment at 927 °C, but decomposition occurred during the cooling process from this temperature; 앫 Decomposition occurred during the 4 min ageing at 510 °C. Although Al was not determined as non-random distribution from FD analysis, the contingency tables of Ni and Al show that these two elements are correlated, Table 4. Contingency table analysis also shows the rejection between Cr and Ni, and between Fe and Cr. No correlation between Cr and Al was detected. It is not surprising to see that Fe and Cr tend to reject each other as Fe is the major element in the matrix and Cr is found to be non-

be those absorbed to the tip surface from the vacuum chamber. 3.1. Ageing at 510 °C 3.1.1. 4 minutes The distribution of the elements seems to be homogeneous from the atom maps of Al, Cr and Ni under this condition, Fig. 4, using Al and Ni as examples. The overall composition of the specimen measured from PoSAP is given in Table 3, with the minor elements being neglected. Its composition agrees well with the nominal composition of

Fig. 4. Atom maps of Al and Ni in the specimen aged for 4 min at 510 °C (the size of Al atoms was made larger to enhance the display).

tage

1.98±0.03

4h

1.26±0.03

7.63±0.06

12.56±0.16

11.94±0.08

13.18±0.05

12.27±0.03

13.94±0.12

Cr

1.86±0.03

1.58±0.06

1.75±0.03

1.49±0.02

1.75±0.01

1.56±0.04

Mo

7.22±0.06

9.20±0.14

9.00±0.07

7.63±0.04

7.97±0.03

8.57±0.10

Ni

0.135±0.009

0.138±0.018

0.111±0.008

0.101±0.004

0.104±0.003

0.118±0.012

Si

81.76±0.09

74.20±0.20

74.89±0.10

75.63±0.06

75.77±0.04

73.74±0.16

Fe

17×17×17

10×9×13

17×16×20

11×11×110

27×26×40

15×14×10

Size (nm3)∗

∗ This is the dimension of the rectangular box enclosing the analysed volume. As the analysed section area reduces as the tip becomes more blunt during a long experiment, the actual analysed volume is smaller than this dimension, albeit usually just slightly.

30 min

2.25±0.07

1.85±0.02

40 min

593 °C 6 min

1.92±0.01

1.93±0.05

Al

15 min

510 °C 4 min

Tage

Table 3 Overall composition of the specimens measured from PoSAP (at%)

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Z. Guo et al. / Acta Materialia 51 (2003) 101–116

Table 4 Contingency tables between Ni (across) and Al (down) in the sample aged for 4 min at 510 °C (atoms per block: 10) Observed table:

0 1 2–10

0 4473 641 51

Expected table: 1 2963 374 29

2–10 1009 101 6

0 1 2–10

0 4521 597 45

1 2946 389 29

2–10 974 128 9

Chi probability: 1% (χ2: 13.2, degrees of freedom (DoF): 4).

randomly distributed. The rejection between Cr and Ni has been reported before in Fe-Cr-Ni ternary alloys [19]. 3.1.2. 15 minutes The atom map of Al in the sample does not clearly show any non-random distribution, Fig. 5, neither do the atom maps of other elements such as Ni, Cr, Mo and Fe. The overall composition of the specimen measured from PoSAP is given in Table 3, which agrees well with the nominal composition of the as-received alloy. FD analysis indi-

Fig. 5.

Al Atom map of sample aged for 15 min at 510 °C.

107

cates that Al, Cr, Fe, and Ni are non-randomly distributed, which can also be revealed by composition profiles. Contingency table analysis was used to test the correlations among different species of atoms. The correlations between Ni and Al, between Cr and Ni, and between Fe and Cr were found to be similar to those in the samples aged for 4 min. Rejection between Cr and Al was also found in this condition. 3.1.3. 40 minutes Detectable particles form after ageing for 40 min at 510 °C. The atom maps of A1 and Ni are shown in Figs. 6(a) and 6(b). Smooth data visualisation (SDV) was carried out to reveal the shape of the precipitates clearly. A two-point-simple smoothing method was used since the 3-dimensional (3D) morphology obtained through this method matches well with the visual impression from atom maps through rotating the analysed volume. It can be seen that Ni-enriched areas and Al-enriched areas are at about the same locations in the analysed box, indicating co-segregation, Figs. 6(c) and 6(d). The iso-surface values of Al and Ni were set as 7.5 and 18 at%, respectively. It should be mentioned that the value of the iso-surface may significantly affect the apparent size of the precipitates in the SDV image. One value was used for each element type throughout this paper to make the particle size somewhat comparable. The composition of the particles is given in Table 5, and the matrix composition is given in Table 6. It can be seen that the particles are enriched with Al and Ni, and depleted of Fe and Cr. Correspondingly, the matrix is depleted of Al and Ni. The difference in the amount of Cr in the matrix and its nominal level is not significant, and the amount of Fe in the matrix is expectedly higher than that in the overall composition. A particle analyses method has recently been developed to quantify particles from 3D atom probe data [30]. The determination of solute-rich regions is performed by connecting solute atoms which lie within a given distance, 0.5 nm in the present case. Then only particles containing above a certain minimum number of solute atoms (ten in this study) are kept for further quantification. Data selected in this way can be used for subsequent

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Fig. 6. Atom maps of Al (a) and Ni (b) after ageing at 510 °C for 40 minutes, and iso-surface maps with iso-surface values of 7.5% for Al (c), and 18% for Ni (d), respectively.

Table 5 Compositions of the particles measured from PoSAP data (at%) Tage

tage

Al

Cr

Ni

Fe

Number of ions

510°C

40 min 4 hr

593°C

6 min

14±2 17±2 26±3 17±2 19±2 16±2 26±2

6±1 6±1 5±1 5±1 5±1 7±1 3±1

18±2 26±3 23±3 24±2 27±2 19±2 35±2

61±2 50±3 44±3 51±3 48±3 57±3 34±2

19±2

4±1

22±2

53±2

2×2×2 nm3 515 ions 1.5×2×1.5 nm3 304 ions 2×2×1.5 nm3 244 ions 1.5×2×2 nm3 351 ions 2.5×1.5×1.3 nm3 331 ions 2×1.5×2 nm3 326 ions 2×2×2 nm3 524 ions (larger precipitate) 2×2×2 nm3 631 ions (smaller precipitate)

30 min

Note: Molybdenum has been found to be randomly distributed throughout the analysed volumes for all treatments. A molybdenum level of around 1 at% is found in all particles

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Table 6 Compositions of the matrix measured from PoSAP data (at%) Tage

tage

510°C 40 min 4h 593°C 6 min 30 min

Al

Cr

Mo

Ni

Fe

Number of ions

0.9±0.2 0.3±0.1 0.4±0.1 0.58±0.03 0.6±0.1

12.7±0.6 12.9±0.5 14±1 7.98±0.12 8.1±0.4

1.4±0.2 1.4±0.2 1.5±0.3 1.87±0.06 2.0±0.2

6.2±0.4 5.7±0.3 6.7±0.6 6.55±0.11 7.2±0.4

78.6±0.8 79.4±0.6 77.5±0.9 82.82±0.16 81.9±0.5

10×2.5×5 nm3 3024 ions 4×14×3 nm3 5349 ions 3×8×2 nm3 1952 ions 16×10×10 nm3 54583 ions 5×5×5 nm3 5210 ions

determination of parameters such as size, shape, solute composition, number density and volume fraction. In the present material, particle analysis shows that in the selected 8.27 × 10⫺24 m3 volume analyses, there are 215 particles containing ten or more Al atoms, 54 particles containing 30 or more, but only 11 particles containing 50 or more Al atoms. It should be noted that the particles of aluminium being referred to here do not necessarily contain 100% aluminium. However, within a particle, at least two aluminium atoms are always spaced within 0.5 nm. The distribution of particle size against the number of particles is shown in Fig. 7. This distribution shows a higher number of small particles (less than 20 Al solute atoms) than advanced particles. This reflects the early stage of particle formation after 40 min at 510 °C. 3.1.4. 4 hours The atom maps of Al and Ni in the sample aged at 510 °C for 4 h are shown in Fig. 8. Since it is difficult to visualise the shape of the precipitates from one direction, the atom maps of Al and Ni recorded from two directions are shown. Combi-

nation of these pictures reveals that the precipitates are of an irregular plate morphology rather than needle-like shape. Smooth data visualisation was also carried out to show the plate morphology. The size of the plate is estimated at about 8–10 nm in diameter and 2 nm in thickness, Fig. 8(c). The composition profiles of Al, Ni, Cr and Fe of a 2 × 2 × 18 nm3 tube are given in Fig. 9. This tube was chosen in a direction perpendicular to the plate surface. The plate thickness estimated from the composition profile of Al and Ni is about 2 nm. It should be noted that the precipitate shape observed from atom maps may somehow deviate from the actual shape of the precipitate due to the evaporation sequence and irregularities of atoms from the specimen and the subsequent reconstruction. The composition of three precipitate particles calculated around their cores are given in Table 5. They are given individually to demonstrate the diversity in the apparent particle composition. Again, these particles are NiAl-enriched and Fe, Cr-depleted. 3.2. Ageing at 593 °C 3.2.1. 6 minutes The atom maps of Al and Ni at this condition are shown in Fig. 10. Smooth data visualisation was performed to show the 3D morphology. The precipitates are found to be needle-like and along the same direction. They are about 2 nm in diameter and 4–5 nm in length. The compositions of two particles and the matrix are listed in Tables 5 and 6, respectively.

Fig. 7. Distribution of Al clusters in the sample aged for 40 min at 510 °C.

3.2.2. 30 minutes The atom maps of Al and Ni in the analysed volume of the sample are shown in Fig. 11,

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Fig. 8. Distribution of of Al and Ni after 4 h ageing at 510 °C. (a), (b) are atom maps from two different directions of projection. Their relationship can be discerned from the different lengths of the edges. (c) is an iso-surface map of Al with isosurface value 7.5%.

together with the figures after SDV treatment. From the SDV figure of Al, it seems that the particle in the middle of the view is needle-like. However, it is not clear what shape the precipitate takes from the SDV of Ni atoms. While viewing the SDV of Al and Ni from different directions, it is interesting to observe that the large NiAl-enriched particle is just a thin layer, Fig. 12. The compositions of the particles and the matrix are again given in Tables 5 and 6. The composition of the overall analysed volume is shown in Table 3. It is unclear why the amount of Cr is low, as all the other samples have compositions close to that of the

Fig. 9. Composition profile of a 2 × 2 × 18 nm3 tube of the sample aged for 4 h at 510 °C (block size: 0.2 nm).

as-received material. It is possible that reverted austenite formed after 30 min at 593 °C. However, the concentrations of Cr in austenite and martensite do not differ much, with marginally less amount in austenite. Thermodynamic calculation using Thermo-Calc gives BCC (55.3 mol%): 80.0Fe-14.2Cr-2.9Ni-1.3 Mo-1.5Al-0.1Ti, and FCC (44.7 mol%): 69.9Fe-12. 1Cr-14.2Ni-1.2Mo-2.6Al-0.03Ti. It seems that even the formation of reverted austenite may not be able to explain the low level of Cr in the specimen aged

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111

Fig. 11. Distribution of Al and Ni in the sample aged for 30 min at 593 °C. Fig. 10. Distribution of Al and Ni in the sample aged for 6 min at 593°C, (a) Atom maps; (b) Iso-surface.

for 30 min at 593 °C. Given that the analysed volume is 17 × 17 × 17 nm3, it is possible that the low level of Cr is due to inhomogeneity of the specimen, since the good accuracy of atom probe analysis of Cr is well established.

4. Discussion 4.1. Microstructure evolution during ageing The amount of retained austenite formed after solution treatment was studied before [7]. The relative amounts of austenite (in wt%) in PH13-8 stainless steel were determined by Mossbauer spectroscopy. After homogenisation at 1121 °C for 4 h (furnace-cooled), then solution treatment at 927 °C for 1.5 h (fan-cooled), with no refrigeration treatment, about 0.4 ± 0.5% retained austenite was detected. Seetharaman et al. solution-treated PH138 at 900 °C for 30 min and their x-ray diffraction results confirmed that there is a complete marten-

Fig. 12. Iso-surface of Al and Ni in the sample aged for 30 min at 593 °C, viewed from a different direction from Fig. 11(b).

site transformation [6]. They estimated the Ms and Mf temperatures to be 60 and 20 °C, respectively. Although in the current work the solution treatment is 1.5 h at 927 °C, it may be reasonable to neglect the influence from retained austenite. The following discussions focused on the evolution of precipitation process during ageing.

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4.1.1. Precipitate type, composition and shape The precipitation progress during ageing at two temperatures was characterised. It is now clear that the particles formed are NiAl-enriched zones. When ageing at 510 °C, such zones start to form at a time between 15 and 40 min. At 593 °C, such precipitates form in less than 6 min ageing. The composition of such zones is far from that of the stoichiometric NiAl phase. However, as can be seen from Table 5 the amounts of Al and Ni in precipitates increase and the amounts of Fe and Cr decrease when ageing prolongs. It was observed in Fe-Ni-Al-Mo systems after prolonged ageing at high temperatures that NiAl precipitates contain only a small amount of Fe [17]. It is known that stoichiometric NiAl exhibits an ordered B2 structure up to its congruent melting point, and the superlattice is stable over a composition range from 41.5 to 55 at% Al at room temperature [31]. It is not clear whether the NiAlenriched phase observed using the atom probe here is of B2 structure since the total of Ni and Al atoms is less than half of the particle composition. The observed superlattice spots by SADP analysis in previous work are in samples after much more prolonged ageing. At an early stage of precipitation, 40 min at 510 °C, the NiAl-enriched particles seem to be spherical. However, when ageing proceeds, the precipitates become platelike (4 h at 510 °C). When ageing temperature is 593 °C, the precipitates formed after 6 min are of needle shape. This is in agreement with the general trend of precipitate sequence in that ‘needles are observed at small supersaturations and plate morphology is observed at large supersaturations’ [32]. 4.1.2. Precipitate size, fraction and number density An obvious increase in precipitate size is observed with increasing ageing at both temperatures. However, the increase in the average particle size cannot be accurately calculated due to the small numbers of particles analysed, neither can the evolution in precipitation fraction. The size of the precipitates in the sample aged for 40 min at 510 °C is spherical and about 1–2 nm in diameter. In the sample aged for 4 h at this temperature, the

plate precipitates are about 2 nm in thickness and 8–10 nm in diameter. After ageing at 593 °C for 6 min, the needle-like precipitates are about 4–5 nm in length and 2 nm in diameter. The number density (Nv) as functions of time and temperature can be estimated in some of the ageing conditions. At 510 °C, detectable precipitates can only be observed after ageing for 40 min. The number density of the Al particles containing ten or more A1 atoms under this condition is estimated at about 2.6 × 1025 / m3. A similar density value should result if the particles are counted using the atom maps or the iso-surface plots in Fig. 6. After 4 h ageing, the particle density decreases to about 2.2 × 1024 / m3 (12 particles in the analysed area 17 × 16 × 20 nm3). Such decrease in number density can only suggest particle coarsening during ageing. 4.1.3. Other precipitates formed during ageing Longer ageing times may allow further development of the composition of the precipitates or new types of precipitates. Thermodynamic calculation with Thermo-Calc shows that Laves phase (Fe,Cr)2Mo may exist in the equilibrium state of PH13-8 steel at 510 or 593 °C. Hochanadel et al. claimed that Laves phase was observed in the cast PH13-8 alloy aged at 621 °C [7]. They also claimed the existence of M23C6 in H1150M condition, with no mention of the methods used to determine its structure. 4.2. Mo segregation at precipitate/matrix interface In the literature, the reported effect of Mo segregation at precipitate/matrix interface differs from alloy to alloy. Stiller et al. found a clear Mo segregation at the interface in 1RK91 maraging steel [23], whereas Miller and Hetherington reported no Mo segregation in the studied Fe-Ni-Al-Mo systems [17]. The current study shows no clear indication that Mo segregates to the precipitate/matrix interface. One of the important features of the studied PH13-8 alloy is that the precipitates are highly resistant to coarsening, which enables the alloy to maintain its mechanical properties during service

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at elevated temperatures. The mean size of the particles is only about 7 nm even after ageing at 625 °C for 4 h in wrought grades [6], and no larger than 50 nm after H1150M treatment of a cast product [7]. The coarsening rate is dependent on the lattice parameter mismatch between the precipitates and the matrix. The Mo-segregation first proposed by Calderon et al. [33] through TEM study on Fe-Ni-Al-Mo systems was disputed by Miller and Hetherington based on atom probe work on the same material [17]. The mechanism in 1RK91 that Mo segregation at particle/matrix interfaces impedes the precipitate coarsening may not apply in PH13-8 alloy. Some of the coarsening theories do predict very sluggish size increment with size proportional to time1/n, where n can be 4 or 5 [34]. This gives much slower coarsening process than the classical relationship with size proportional to time1/3. In addition, the coarsening process of precipitates can be accompanied with the formation of reverted austenite during ageing. Reverted austenite is a phase which contains much higher Ni and Al contents than the matrix. Its formation will undoubtedly decrease the supply of Ni and Al atoms to the NiAl-enriched precipitates, as, effectively, some Ni and Al atoms dissolved from the smaller precipitates participate in the formation of austenite. 4.3. Spinodal decomposition of chromium Cr is tested to be of non-random distribution from binomial FD analysis in all the ageing conditions. Two methods were applied to measure the spinodal amplitude of Cr, both based on the comparison of the sample and a model frequency distribution. One is Cahn’s sinusoidal or linear Pa model, and the other Langer, Bar-on and Miller (LBM) or non-linear model. In each case, the model distribution is compared with the measured distribution using the method of maximum likelihood and χ2 estimator is used to find the degree of fit. Advantages and disadvantages of these two models were discussed by Hetherington et al. [35]. Results from the two models are given in Table 7. The Chi probability values are in all but one case 100% which means that both models for spinodal decomposition fit well with the observed data.

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However, it should be pointed out that the measured amplitudes, Pa or µ2 ⫺ µ1, are comparable to the errors in the average composition in one block of 100 atoms, √c(1 ⫺ c) / N (3%). Nevertheless, such analysis shows some indication of a spinodal decomposition of Cr elements. Analysis using these two models was also carried out on other elements that show non-random distribution, Fe, Ni, and Al. Their distributions show significant deviation from the two models and therefore their decompositions are not of spinodal type. Chromium redistribution in low-Cr steels was observed in an aged ferritic-martensitic steel (with 11 at% Cr) using atom probe [36]. Due to the high carbon content (0.46–0.83 at%), the actual Cr in the matrix is even lower since Cr takes part in forming M23C6 carbides. After ageing at 400 °C for 17000 h, long range fluctuations, with concentration difference of 2.9 at% were found. Therefore, in low-Cr steels, the decomposition of Cr can still follow a spinodal pattern. The possible relation between such spinodal decomposition process and the hardening effects is discussed in the following section. 4.4. Hardening mechanisms No detectable precipitates have formed in samples aged for 4 and 15 min at 510 °C, though the hardening effects at these ageing conditions correspond to about 20% and over 40% of the total hardness increase, respectively (Fig. 1). Similar phenomenon was observed for PH17-4 alloy by Murayama et al. [22]. The hardening effect was concluded to be caused by the Cr spinodal decomposition since they did not observe any precipitates when significant hardening was evident. However, from the PoSAP study, it is impossible to determine which element starts to re-distribute first (spinodal decomposition, in terms of redistribution, however, assumes Cr is the first element to start to decompose). The hardening effects observed during early ageing should be due to the redistribution of atoms such as Fe, Cr, Ni, and/or Al. While studying another precipitation hardening system Al-Mg-Cu, Reich et al. [37] suggested that the initial rapid hardening during ageing is associated with a solute-dislocation interaction, i.e. the solute segregation to the existing dislocations

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Table 7 Pa and LBM analysis of the decomposition of Cr (atoms per block: 100) Tage

tage

510 °C

4 min 15 min 40 min 4h 6 min 30 min 1h

593 °C

χ2/DoF

Pa 2.2 1.7 2.7 2.9 2.5 2.2 1.8

20.6/16 28.1/20 29.2/22 35.2/19 11.1/15 11.6/16 9.1/10

Prob. (%) 100 100 100 1.3 100 100 100

which causes dislocation locking. The zones formed through atom redistribution or the solute-dislocation interaction caused by solute segregation to dislocations are thought to contribute to the early hardening effects. The correlation between Al and Ni changes during the ageing process. In the samples aged for 4 and 15 min at 510 °C where no precipitates are observed using atom probe, contingency tables, Table 4, show that they reject each other. In the other studied samples, there is a clear indication of co-segregation of Ni and Al, Table 8. It is possible that Al or Ni atom tends to form vacancy-atom complexes, which are then driven to vacancy sinks, i.e. dislocations, resulting in the apparent rejection between Ni and Al. When ageing proceeds, the vacancy sinks become fully occupied, and the affinity between Ni and Al starts to operate, resulting in the co-segregation of Ni and Al as observed. Strong indication of the formation of vacancy-Mg complex was shown in an Al-Cu-Mg alloy [38]. Table 8 Contingency tables between Ni (across) and Al (down) in the sample aged for 40 min at 510 °C (atoms per block: 10) Observed table:

0 1 2 3 4–10

0 34673 2876 425 88 35

Expected table: 1 17213 2004 397 112 31

2–10 6587 1177 263 41 14

0 1 2 3 4–10

0 33784 3499 626 139 45

1 17520 814 325 72 22

2–10 7164 740 129 28 7

Chi probability: 0.1% (χ2: 743.7, degrees of freedom (DoF): 8).

µ1

µ2

s

0.5 0.5 1.5 1.0 1.0 1.5 0.5

3.5 2.5 2.5 3.5 3.0 1.5 3.5

1.0 0.5 0.5 1.0 0.5 0.5 0.5

χ2/DoF 18.0/17 21.9/20 28.8/21 14.4/19 10.0/15 12.4/16 7.6/10

Prob. (%) 100 100 100 100 100 100 100

Modelling the precipitation hardening kinetics has been a long term interest. It should be noted that precipitate hardening theories usually assume the composition of the precipitate to be stoichiometric and then correlate the strength change with the precipitate size and fraction. In a more complete, physical model, the composition change during ageing should be taken into account. The growth/coarsening kinetics should be revised accordingly as well.

5. Conclusions The current study of the precipitation process during ageing in a wrought PH13-8 steel has led to the following conclusions. 1. Detectable precipitates form after 40 min ageing at 510 °C, or 6 min at 593 °C. They are enriched in Ni and Al, but depleted of Fe and Cr. The amount of Ni and Al increases when ageing proceeds, but the precipitate composition is far from the stoichiometric NiAl B2 phase after 4 h at 510 °C (H950), or 30 min at 593 °C. 2. Ageing for 40 min at 510 °C results in spherical NiAl-enriched particles. The precipitates may take on different shapes at different ageing temperatures. 3. Coarsening processes take place simultaneously with the development of the composition of the NiAl-enriched precipitates. 4. No convincing sign of Mo segregation at the precipitate/matrix interface was found. 5. The decomposition of Cr differs from that of

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the other elements. It is possible that spinodal decomposition of Cr takes place in this lowCr steel. 6. The hardening effects observed during early stages of ageing should be due to the redistribution of atoms, possibly through the solute-dislocation interaction caused by solute segregation to dislocations.

Acknowledgements G.D.W. Smith and A. Cerezo in the Department of Materials at Oxford University are thanked for providing the position sensitive atom probe facility and software for data analysis, and useful discussions. Thanks are also given to T.J. Godfrey and other personnel working in the atom probe laboratory. This work was carried out within the project of ‘Computer Modelling of the Evolution of Microstructure during Processing of Maraging Steels’ sponsored by the Engineering and Physical Sciences Research Council, UK, under Grant No. GR/N08971.

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