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A films at grain boundaries on crack propagation in an ultra-fine acicular ferrite pipeline steel
Acta Materialia 54 (2006) 435–443 www.actamat-journals.com
In situ TEM study of the effect of M/A films at grain boundaries on crack propagation in an ultra-fine acicular ferrite pipeline steel Yong Zhong
a,*
, Furen Xiao b, Jingwu Zhang b, Yiyin Shan a, Wei Wang a, Ke Yang
a
a
b
Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China College of Materials Science and Engineering, Yanshan University, Qinhuangdao 066004, China
Received 27 May 2005; received in revised form 15 September 2005; accepted 15 September 2005 Available online 2 November 2005
Abstract Microstructural refinement of structural materials generally improves their tensile properties but deteriorates their fatigue properties. However, pipeline steels with ultra-fine acicular ferrite (UFAF) possess not only high strength and toughness, but also a low fatiguecrack-growth rate (FCGR) and long fatigue-propagation life. In this paper, the micro-fracture mechanisms of an UFAF pipeline steel are investigated by in situ tensile testing in a transmission electron microscope. The results indicate that a grain-boundary-film structure composed of martensite/austenite could significantly influence the crack propagating behavior in the UFAF steel, consequently lowering the FCGR by enhancing roughness-induced crack closure during cyclic loading. 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: In situ TEM; Fatigue crack growth; Grain boundary structure; Crack closure; Pipeline steels
1. Introduction One of the major concerns in maintaining pipeline integrity is severe damage such as leakage or fracture resulting from defects in the pipes during operation. Since most pipelines are used under working conditions involving inner pressure cycling or frequent stops and starts, they are usually subjected to fatigue loading [1]. In previous work by the authors, the effect of microstructure on the fatigue behavior of low-carbon micro-alloyed pipeline steels was investigated [2]. The fatigue-crackgrowth rates (FCGR) and fatigue-propagation lives under simulated operating conditions were determined for pipeline steels with various microstructures and grain sizes, including polygonal ferrite plus pearlite (PF+P), acicular ferrite (AF), and ultra-fine acicular ferrite (UFAF). The results indicated that, under the experimental conditions employed, the UFAF possessed the lowest FCGR and the
*
Corresponding author. Tel.: +81 82 424 5744. E-mail address: [email protected] (Y. Zhong).
greatest fatigue life, as well as the best mechanical properties, including high strength and excellent toughness. This observation provided considerable support for the application of UFAF steels for pipelines. However, the results were not easily understood, since a refined microstructure usually leads to good tensile properties but poor fatigue properties, such as a high FCGR and low fatigue threshold [3]. The FCGR can be lowered markedly by a plastic zone at the crack tip leaving residually stretched material in the wake of crack, causing a premature interference between the crack surfaces [4], or so-called Elber closure [5]. Elber closure can be induced by many factors such as plasticity, roughness of crack surfaces, oxide, reinforcing particles, and so on. In pipeline steels, one of the most important factors that can cause Elber closure is the roughness-induced crack closure (RICC). The large grain size leads to increases in heterogeneity of deformation, and this heterogeneity leads to better slip reversibility [6] and larger crystallographic facets on the fracture surfaces that lead to more RICC. Furthermore, the crack paths become more tortuous and deviate from the Mode I plane [7,8]. All these effects combine to produce
1359-6454/$30.00 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2005.09.015
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Y. Zhong et al. / Acta Materialia 54 (2006) 435–443
higher intrinsic resistance against crack growth in a coarsegrained material [9,10]. Considering the inconsistency of the results obtained with the RICC mechanism, it may be concluded that there must be some special mechanisms that could have a strong influence on the RICC and consequently the fatigue growth behavior in a UFAF pipeline steel. The aim of this paper is to investigate the fatigue mechanism in a UFAF pipeline steel by conducting in situ tensile tests in a transmission electron microscope (TEM) during which the deformation behavior and crack propagation in the UFAF can be observed directly. In pipeline steel technology, comparatively little research has been conducted on the merging of fatigue behavior and fracture mechanisms. These results would provide industry with a unified approach to the improvement of the reliability and economic usefulness of oil and gas transmission pipelines. 2. Experimental materials and procedures The specimens for in situ TEM observation came from the hot rolled plates of an UFAF pipeline steel. After being
mechanically processed to rectangular foil with the dimensions 10 mm long, 3 mm wide and 30 lm thick, the specimens were electro-polished by a twin-jet electropolisher in a solution of 10% perchloric acid and 90% acetic acid. The resulting specimen geometry is shown schematically in Fig. 1. The specimens were then strained on a tensile stage in a Hitachi H800 (200 kV) TEM. The deformation and the propagation of microcracks in the steel were thus observed in situ. To investigate the effects of microstructures on fatigue mechanism, two other pipeline steels were used to compare the fatigue crack propagation behaviors by optical observation. One of them is a commercial API-X60 pipeline steel and the other is an AF pipeline steel with the same chemical composition as the sample UFAF, but with a coarser grain size. The chemical compositions and mechanical properties of the samples are shown in Tables 1 and 2, respectively. For optical observation, specimens were mechanically polished and etched in 3% nital solution, and the microstructures are shown in Fig. 2. Both UFAF and AF were produced by a 7-step controlled hot rolling. After hot rolling, the plates were
Fig. 1. Mechanism of in situ TEM tensile specimen. Table 1 Chemical compositions of the experimental steel (elements in wt.% except P, S, O and N in ppm) Steels
C
Si
Mn
Mo
Ni
Cu
Ti
Nb
V
P
UFAF AF
0.025
0.24
1.56
0.32
–
–
–
0.039
0.019
20
API-X60
0.07
0.25
0.9
–
0.2
0.2
0.015
0.04
0.04
70
S
O
N
6
43
62
15
40
40
Table 2 The mechanical properties of the samples Samples
Ultimate strength (MPa)
Yield strength (MPa)
Elongation (%)
Charpy-V-notched impact energya (J)
UFAF AF API-X60
650 679 519
587.5 577.5 454
24 22 29
148 112 64
a
Half size specimens (5 · 10 · 55 mm), at 40 C.
Fig. 2. The microstructures of the samples: (a) UFAF; (b) AF; (c) X60.
Y. Zhong et al. / Acta Materialia 54 (2006) 435–443
437
Table 3 The control rolling and accelerated cooling processing of UFAF and AF steel Steel
Start-rolling temperature(C)
Final rolling temperature(C)
Cooling rate (C)
Coiling temperature (C)
UFAF AF
1056 1070
800 800
35 21.44
600 600
water-cooled down to around 600 C immediately, and then held at 600 C for 1 h to simulate the coiling process of steel plates, and finally furnace cooled to room temperature. Some processing parameters were finely adjusted to obtain various grain size and properties. As shown in Table 3, the AF steel had a higher start-rolling temperature, and lower cooling rate, and consequently a larger grain size compared to the UFAF. In order to quantitatively analyze the effect of RICC, the crack-surface roughness of the FCGR sample of steel X60 and UFAF was measured on a 2205-type multi-parameters profilometer with a diamond stylus of 1 lm radius, in accordance with the standard method of ISO 468-1982. Table 4 The crack surface roughness of X60 and UFAF Steel
R (lm)
S (lm)
H ()
X60 UFAF
8.2 24.4
39.7 63.8
17.9 40.6
The measuring range of the profilomater is 50 lm and the resolution is 0.1 lm. The mean height of profile irregularities (R), the mean spacing of the profile irregularities (S) and the mean slope of the profile (H) were statistically measured. In every sample, five sampling locations were measured with the traverse length of the stylus of 2 mm and the traverse speed of 1 mm/s in each sampling location. The final results were averaged from these five sampling locations, as given in Table 4. 3. Results 3.1. Effect of a high angle grain boundary (HAGB) on crack propagation In Fig. 3(a), there is a high-angle grain boundary (HAGB) between grain A and grain B. The loading axis is along the vertical direction. A crack propagated in grain A from right to left in a zigzag fashion. As the sample was
Fig. 3. Crack propagating through a high angle grain boundary (HAGB): (a) a crack propagated from grain A to grain B; (b) the dislocations in grain B moved under loading; (c) the crack passed through HAGB and changed its propagating direction; (d) the crack propagated in grain A along the slip direction of dislocation, and a clear dislocation free zone (DFZ) formed in front of the crack tip.
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strained, dislocations in grain A moved and piled up at front of the HAGB. When loading was increased, the dislocations in grain B started to slip (Fig. 3(b)). It can be observed that the slip direction in grain B was different from that in grain A, and the misorientation between A and B was about 30. With further increase in loading, the crack propagated across the HAGB with a significant deflection of its trajectory. In addition, a dislocation free zone (DFZ) was formed at the tip of the crack (Fig. 3(c)). After propagating in grain B for some distance, the propagation direction was gradually changed to follow the direction of the DFZ (Fig. 3(d)). 3.2. Effect of low-angle grain boundary (LAGB) on crack propagation Fig. 4 showed the process of crack propagation across a low-angle grain boundary (LAGB). In this case, grains A and B were separated by a LAGB. The loading axis was along the vertical direction. Under loading, a crack propagated from the upper right corner towards the lower left in grain A (Fig. 4(a)). As the loading increased, dislocations were generated at the LAGB in grain B, and slipped towards the left side. It was observed that the dislocation density in this grain was far higher at the higher strain
(Fig. 4(b)) than that at the lower strain (Fig. 4(a)). When the loading increased further, the crack propagated close to the LAGB, and the dislocations kept being generated from the LAGB. Some dislocations that piled up at the LAGB in grain A seemed to slip over the grain boundary to grain B. It would be noticed that the crack propagation in grain A and dislocation glide in grain B were going in almost the same direction (Fig. 4(c)). Finally, the crack propagated across the LAGB suddenly, and the propagation direction of crack in grain A and grain B did not change perceptibly (Fig. 4(d)). 3.3. Effect of LAGB with an M/A film structure on crack propagation Fig. 5 shows a crack propagating across a special LAGB. The loading axis was still along the vertical direction. Grains A and B were separated by a LAGB. It should be noticed that the morphological character of the LAGB is quite different from that in Fig. 4. The thickness of this LAGB was larger than that of a normal LAGB, and some fine structure can be observed (Fig. 5(a)). As per the previous report of the authors on microstructure characteristics of acicular ferrite [11], this kind of structure was proved to be a continuous film with thickness of several to tens of nanometers at the boundary between
Fig. 4. Crack propagating across low angle grain boundary: (a) low angle grain boundary between grains A and B; (b) crack propagating and boundary emitting dislocation; (c) crack approaching to the boundary; (d) crack propagating across the boundary.
Y. Zhong et al. / Acta Materialia 54 (2006) 435–443
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Fig. 5. Crack propagating across the low angle grain boundary with M/A film structure: (a) low angle grain boundary between grains A and B; (b) crack approach to boundary; (c) crack deflecting at the boundary; (d) crack propagating across the boundary.
acicular ferrite grains (Figs. 6(a) and (b)). For determining the structure of the film, TEM micro-diffraction and X-ray diffraction (XRD) were adopted. The results showed that the thin film is mainly a body-centered cubic (bcc) structure, which in steel is well known to be either ferrite or martensite (Fig. 6(c)). Therefore, this thin film should be composed of martensite and retained austenite (M/A), which transformed from retained austenite during the simulated coiling process after accelerated cooling. This M/A film in the UFAF pipeline steels, somewhat similar to the thin film reported in other conventional microstructures such as the welded AF microstructure or bainite [12,13], is significantly different in morphology from the second phase in the AF conventional carbon pipeline steels, in which the M/A component is distributed as islands. Since no visual result of the effect of M/A film structure on defor-
mation and crack propagation was reported before, it would be of great interest to investigate the process of crack propagation across such a structure. In Fig. 5(a), a crack was seen to propagate in grain A from the right side to left under loading. With increased loading, the crack propagated close to the LAGB, but no significant change occurred in the dislocation distribution in either grain A or grain B (Fig. 5(b)). When loading was further increased, the crack propagated across the LAGB, and the propagation direction of the crack deflected markedly upon crossing the LAGB (Fig. 5(c)). After the crack had propagated some distance in grain B, it was obvious that the misorientation of crack propagation between grains A and B was close to 90 (Fig. 5(d)). According to crystallographic analysis of dislocation movement and crack propagation, the slip direction in
Fig. 6. Film structure at the grain boundary of acicular ferrite: (a) bright field of M/A film structure at the grain boundary; (b) dark field of M/A film structure at the grain boundary; (c) micro-diffraction pattern of M/A film.
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bcc crystals is Æ1 1 1æ, which has fourfold symmetry [14]. Therefore, the misorientation of the slip direction up to 90 may be just the different slip direction of the same slip system. 4. Discussion 4.1. Mechanism of RICC Our previous work showed that the FCGR of UFAF steel in the Paris regime [15] was significantly less than that of X60, and also a little bit lower than that of the steel with an AF microstructure [2]. For example, at DK = 35 MPa m0.5, the FCGR of UFAF was almost half of that of X60, as shown in Fig. 7. To investigate the mechanism for the resistance against fatigue crack propagation of UFAF, trajectories of fatigue crack propagation of various samples were observed, as shown in Fig. 8. It can be seen that the crack in the X60 sample did not show much deviation and remained quite close to a straight crack. On the other hand, the crack in
da/dN, mm/cycle
X60 AF UFAF
1E-3
1E-4
30
40
∆K, MPa.m
50
60
70
80
90
the UFAF steel kinked and meandered significantly, which could lead to a rougher crack surface with more asperities. The deflection of the crack in the AF sample was intermediate between that in the X60 and UFAF samples. Therefore, it is further suggested there are some special mechanisms in the UFAF that can affect the crack propagation behavior by deflecting the crack. According to the results of the in situ TEM observations, the effect of a HAGB on deformation and crack propagation was strongly different from that of a LAGB. The change in operating slip system, which is necessary as a crack crosses a grain boundary, results in the grain boundaries being barriers for dislocation motion. For a HAGB, the formation of a DFZ in a neighboring grain will depend on the operation of its own slip system, which will lead to the significant difference of direction and crack length between the two sides of the grain boundary. Thus the crack propagation direction will change obviously after crossing the grain boundary, and the larger the misorientation between two grains, the more pronounced will be the deceleration in crack-growth rate at the grain boundary [16]. Conversely, for an LAGB, the small difference in orientation between neighboring grains would produce a relatively small retardation. Brown [17] also claimed that an LAGB fails to provide a significant barrier to crack propagation. However, the observations from in situ TEM indicated that in the UFAF steel, the LAGB with an M/A film had a similar effect to the HAGB on the dislocation motion and crack propagation. It indicated that even at an LAGB, the M/A film can be an effective barrier to dislocation motion and crack propagation, despite the small misorientation between neighboring grains. Consequently, in the UFAF it can be concluded that not only the crystallographic misorientation between two neighboring grains, but also the character of the grain boundary itself affected the crack propagation. The M/A film at grain boundary can increase the possibility of crack deflection, which leads to larger asperities than would occur in a PF+P structure as in X60. 4.2. Estimating the effect of RICC by mathematical models
0.5
Fig. 7. Fatigue crack propagation rates of X60, AF and UFAF.
As described in the above section, the UFAF pipeline steel exhibited better fatigue properties than coarse PF+P
Fig. 8. Trajectories of fatigue crack propagation of various samples: (a) coarse polygonal ferrite plus pearlite; (b) acicular ferrite; (c) ultra-fine acicular ferrite.
Y. Zhong et al. / Acta Materialia 54 (2006) 435–443
441
Fig. 9. Scheme of the crack closure mechanism.
and coarse AF, and this improvement was the result of increased roughness in the crack surface. It is of interest, therefore, to analyze the resistance to fatigue crack propagation of a deflected crack in terms of present mathematical models and hypotheses based on the RICC mechanism. A quantitative assessment of the RICC level is possible by using the statistical modal based on a simple structural analysis [18]. The two-dimensional zigzag crack paths characterized by an average angle h, which is considered to be the determining factor for the direct effect of crack deflection, is assumed to produce the RICC according to the scheme in Fig. 9. During the loading part of the cycle, the crack tip is opened in mixed mode reaching the maximum crack tip opening displacement denoted as dmax at the peak stress (Fig. 9(b)). This displacement can be considered to be composed of both the reversible normal component d1 and the irreversible shear component d2 defining the local mixed mode (I + II). During the unloading phase, consequently, the shear displacement is not recovered so that the crack surfaces come into the contact with each other before the applied stress becomes zero [19], as shown in Fig. 9(c). The crack path deflection reduces the crack growth rate in two ways: (i) it causes a direct reduction of the local driving force for crack propagation and an increase in the total length of crack path [7]; and (ii) it causes a reduction of DK experienced at the crack tip by RICC from the nominally applied DK to the effective DKeff. These two effects are referred to as the extrinsic retardation mechanisms [20]. For the fatigue crack growth at a loading ratio of 0, the
relationship between DK and DKeff can be demonstrated as [13,21]: sffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi) ( v tan h 2 h DK; ð1Þ DK eff ¼ cos 1 2 1 þ v tan h lI 3 Rs 1 ; ¼ lII 2 Rs þ 1 1 ; Rs ¼ cos h
v¼
ð2Þ ð3Þ
where v and Rs are dimensionless parameters representing the mismatch between the crack surfaces and the roughness of the fracture surface, respectively. According to the data of crack surface roughness, the mean slope of crack profile (H), i.e., the average deflection angles h of X60 and UFAF are 17.9 and 40.6 (Table 4), respectively. Substituting the h values into Eq. (1), the effective crack propagation driving force for two microstructures would be: DK effðX60Þ ¼ 0:87DK; DK effðUFAFÞ ¼ 0:54DK. The results are consistent with the experimental data, which indicates that, due to RICC caused by crack deflection during cyclic loading, the DKeff of UFAF is much lower than that of X60. Substituting DKeff for DK in Fig. 7, the normalized fatigue crack growth rates of X60 and UFAF show that the fatigue growth rate of X60 would be lower than that of UFAF if the effect of RICC was eliminated. That is, X60 possesses a high intrinsic resistance
Fig. 10. Trajectories of fatigue crack of (a) X60 and (b) UFAF.
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Y. Zhong et al. / Acta Materialia 54 (2006) 435–443
∆κ eff-da/dN, UFAF ∆κ eff-da/dN, X60
da/dN, mm/cycle
∆κ -da/dN, X60 ∆κ -da/dN, UFAF
1E-3
1E-4 20
30
40
∆K, MPa.m
50
60 70 80 90
0.5
Fig. 11. Fatigue crack growth rate curves of X60 and UFAF vs. DK and DKeff.
against fatigue crack propagation resulting from matrix properties (Fig. 11). Therefore, in spite of the low intrinsic crack growth rates, the high observed crack growth rate can be due to a low level of crack closure. It was observed that, for the X60 sample, the deviation of the crack path from the macroscopic crack plane was about 10 lm, which is close to the PF grain size, while that of the UFAF sample was about 30 lm, which is much larger than the UFAF grain size (Fig. 10). So, it is further confirmed that RICC played a significant role in increasing the resistance against crack growth in the UFAF steel. Based on the analysis above, it is implied that the special microstructure constituent, i.e., the M/A film structure at the grain boundary can have a strong influence on deflecting the crack trajectories and enhancing RICC. Therefore, it is hoped that, in the future, more intensive study will be carried out to further develop the necessary understanding of the nature and role of the grain boundary films, in order to further improve the mechanical behavior of these high strength pipeline steels. 5. Conclusions The micro-fracture behavior of a high strength pipeline steel with a UFAF microstructure was directly observed by in situ TEM tensile testing in order to study the effect of the grain boundary structures on the FCGR. According to the results, the following conclusions can be made.
(1) A LAGB did not influence the crack propagation significantly, but an LAGB with an M/A film structure was an effective barrier to dislocation motion and crack propagation, despite the small misorientation between neighboring grains. Consequently, such boundaries can lead to strongly deflected fatigue crack trajectories, which, in turn, lead to rougher asperities on the crack surfaces during cyclic loading. (2) The mathematical analysis of the fatigue mechanism indicated that despite a poor intrinsic resistance against crack growth, owing to refined grain size, the UFAF possesses a lower effective DK (DKeff) than coarse polygonal ferrite plus pearlite and coarse AF, due to a higher level of RICC that is attributed to the increasing in asperities on the crack surfaces. Therefore, by enhancing RICC, the M/A films at grain boundaries play a significant role in the improvement of the resistance against fatigue crack growth in UFAF. (3) This RICC mechanism enhanced by the M/A film structure, which is a novel discovery for the fatigue resistance mechanism in low carbon micro-alloyed pipeline steels, provides a new idea for the development of pipeline steels with both high fatigue properties and high strength. Acknowledgments The work was financially supported by funding from the National Natural Sciences Foundation (No. 50471106). The authors thank Dr. William Warke from BP Amoco for amending the paper and Dr. Xiaoqiang Zhang from our institute for the help with experiments. References [1] Hagiwara N, Meziere Y, Oguchi N, Zarea M, Champavere R. JSME Int J Ser A 1999;42(4):610. [2] Zhong Y, Shan Y, Xiao F, Yang K. Mater Lett 2005;59(14–15): 1780. [3] Suresh S. Fatigue of materials. 2nd ed. Cambridge: Cambridge University Press; 1998. p. 257. [4] Suresh S, Ritchie RO. Metall Trans A 1982;13:1627. [5] Elber W. Eng Fract Mech 1970;2:37. [6] Lerch BA, Jayaraman N, Antolovich SD. Mater Sci Eng 1984;66:151. [7] Kitagawa H, Yuuki R, Ohira T. Eng Fract Mech 1975;7:515. [8] Campbell JP, Venkateswara Rao KT, Ritchie RO. Metall Mater Trans A 1999;30:563. [9] King JE. Mater Sci Tech 1987;3:750. [10] Vinogradov A, Nagasaki S, Patlan V, Kitagawa K, Kawazoe M. Nanostruct Mater 1999;11(7):925. [11] Zhao MC, Hanamura T, Qiu H, Ke Yang. Mater Sci Eng A 2005;395(1–2):327. [12] Spanos G, Moon DW, Fonda RW, Menon ESK, Fox AG. Metall Mater Trans A 2001;32:3043. [13] Chang LC, Bhadeshia HKDH. Mater Sci Technol 1995;11:874. [14] Hull D, Bacon DJ. Introduction to dislocations. 3rd ed. Oxford: Pergamon Press; 1984. p. 257. [15] Paris PC, Tada H, Donald JK. Int J Fatigue 1999;21:S35. [16] Chen D, Sixta ME, Zhang XF, De Jonghe LC, Ritchie RO. Acta Mater 2000;48:4599.
Y. Zhong et al. / Acta Materialia 54 (2006) 435–443 [17] Brown MW. In: Miller KJ, de los Rios ER, editors. The behavior of short fatigue cracks. London: MEP; 1986. p. 423. [18] Pokluda J, Sandera P, Hornikova J. Fatigue Fract Eng Mater Struct 2004;27:141.
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[19] Wang S-H, Muller C, Exner HE. Metall Mater Trans A 1998;29:1933. [20] Shang JK, Tzou JL, Ritchie RO. Metall Trans A 1987;18:1613. [21] Noda N, Oda K. In: Murakami Y, editor. Stress intensity factors handbook, vol. IV. Kyoto, Japan: TSMS and Elsevier; 2001. p. 402.
In situ TEM study of the effect of M/A films at grain boundaries on crack propagation in an ultra-fine acicular ferrite pipeline steel Yong Zhong
a,*
, Furen Xiao b, Jingwu Zhang b, Yiyin Shan a, Wei Wang a, Ke Yang
a
a
b
Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, China College of Materials Science and Engineering, Yanshan University, Qinhuangdao 066004, China
Received 27 May 2005; received in revised form 15 September 2005; accepted 15 September 2005 Available online 2 November 2005
Abstract Microstructural refinement of structural materials generally improves their tensile properties but deteriorates their fatigue properties. However, pipeline steels with ultra-fine acicular ferrite (UFAF) possess not only high strength and toughness, but also a low fatiguecrack-growth rate (FCGR) and long fatigue-propagation life. In this paper, the micro-fracture mechanisms of an UFAF pipeline steel are investigated by in situ tensile testing in a transmission electron microscope. The results indicate that a grain-boundary-film structure composed of martensite/austenite could significantly influence the crack propagating behavior in the UFAF steel, consequently lowering the FCGR by enhancing roughness-induced crack closure during cyclic loading. 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: In situ TEM; Fatigue crack growth; Grain boundary structure; Crack closure; Pipeline steels
1. Introduction One of the major concerns in maintaining pipeline integrity is severe damage such as leakage or fracture resulting from defects in the pipes during operation. Since most pipelines are used under working conditions involving inner pressure cycling or frequent stops and starts, they are usually subjected to fatigue loading [1]. In previous work by the authors, the effect of microstructure on the fatigue behavior of low-carbon micro-alloyed pipeline steels was investigated [2]. The fatigue-crackgrowth rates (FCGR) and fatigue-propagation lives under simulated operating conditions were determined for pipeline steels with various microstructures and grain sizes, including polygonal ferrite plus pearlite (PF+P), acicular ferrite (AF), and ultra-fine acicular ferrite (UFAF). The results indicated that, under the experimental conditions employed, the UFAF possessed the lowest FCGR and the
*
Corresponding author. Tel.: +81 82 424 5744. E-mail address: [email protected] (Y. Zhong).
greatest fatigue life, as well as the best mechanical properties, including high strength and excellent toughness. This observation provided considerable support for the application of UFAF steels for pipelines. However, the results were not easily understood, since a refined microstructure usually leads to good tensile properties but poor fatigue properties, such as a high FCGR and low fatigue threshold [3]. The FCGR can be lowered markedly by a plastic zone at the crack tip leaving residually stretched material in the wake of crack, causing a premature interference between the crack surfaces [4], or so-called Elber closure [5]. Elber closure can be induced by many factors such as plasticity, roughness of crack surfaces, oxide, reinforcing particles, and so on. In pipeline steels, one of the most important factors that can cause Elber closure is the roughness-induced crack closure (RICC). The large grain size leads to increases in heterogeneity of deformation, and this heterogeneity leads to better slip reversibility [6] and larger crystallographic facets on the fracture surfaces that lead to more RICC. Furthermore, the crack paths become more tortuous and deviate from the Mode I plane [7,8]. All these effects combine to produce
1359-6454/$30.00 2005 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2005.09.015
436
Y. Zhong et al. / Acta Materialia 54 (2006) 435–443
higher intrinsic resistance against crack growth in a coarsegrained material [9,10]. Considering the inconsistency of the results obtained with the RICC mechanism, it may be concluded that there must be some special mechanisms that could have a strong influence on the RICC and consequently the fatigue growth behavior in a UFAF pipeline steel. The aim of this paper is to investigate the fatigue mechanism in a UFAF pipeline steel by conducting in situ tensile tests in a transmission electron microscope (TEM) during which the deformation behavior and crack propagation in the UFAF can be observed directly. In pipeline steel technology, comparatively little research has been conducted on the merging of fatigue behavior and fracture mechanisms. These results would provide industry with a unified approach to the improvement of the reliability and economic usefulness of oil and gas transmission pipelines. 2. Experimental materials and procedures The specimens for in situ TEM observation came from the hot rolled plates of an UFAF pipeline steel. After being
mechanically processed to rectangular foil with the dimensions 10 mm long, 3 mm wide and 30 lm thick, the specimens were electro-polished by a twin-jet electropolisher in a solution of 10% perchloric acid and 90% acetic acid. The resulting specimen geometry is shown schematically in Fig. 1. The specimens were then strained on a tensile stage in a Hitachi H800 (200 kV) TEM. The deformation and the propagation of microcracks in the steel were thus observed in situ. To investigate the effects of microstructures on fatigue mechanism, two other pipeline steels were used to compare the fatigue crack propagation behaviors by optical observation. One of them is a commercial API-X60 pipeline steel and the other is an AF pipeline steel with the same chemical composition as the sample UFAF, but with a coarser grain size. The chemical compositions and mechanical properties of the samples are shown in Tables 1 and 2, respectively. For optical observation, specimens were mechanically polished and etched in 3% nital solution, and the microstructures are shown in Fig. 2. Both UFAF and AF were produced by a 7-step controlled hot rolling. After hot rolling, the plates were
Fig. 1. Mechanism of in situ TEM tensile specimen. Table 1 Chemical compositions of the experimental steel (elements in wt.% except P, S, O and N in ppm) Steels
C
Si
Mn
Mo
Ni
Cu
Ti
Nb
V
P
UFAF AF
0.025
0.24
1.56
0.32
–
–
–
0.039
0.019
20
API-X60
0.07
0.25
0.9
–
0.2
0.2
0.015
0.04
0.04
70
S
O
N
6
43
62
15
40
40
Table 2 The mechanical properties of the samples Samples
Ultimate strength (MPa)
Yield strength (MPa)
Elongation (%)
Charpy-V-notched impact energya (J)
UFAF AF API-X60
650 679 519
587.5 577.5 454
24 22 29
148 112 64
a
Half size specimens (5 · 10 · 55 mm), at 40 C.
Fig. 2. The microstructures of the samples: (a) UFAF; (b) AF; (c) X60.
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Table 3 The control rolling and accelerated cooling processing of UFAF and AF steel Steel
Start-rolling temperature(C)
Final rolling temperature(C)
Cooling rate (C)
Coiling temperature (C)
UFAF AF
1056 1070
800 800
35 21.44
600 600
water-cooled down to around 600 C immediately, and then held at 600 C for 1 h to simulate the coiling process of steel plates, and finally furnace cooled to room temperature. Some processing parameters were finely adjusted to obtain various grain size and properties. As shown in Table 3, the AF steel had a higher start-rolling temperature, and lower cooling rate, and consequently a larger grain size compared to the UFAF. In order to quantitatively analyze the effect of RICC, the crack-surface roughness of the FCGR sample of steel X60 and UFAF was measured on a 2205-type multi-parameters profilometer with a diamond stylus of 1 lm radius, in accordance with the standard method of ISO 468-1982. Table 4 The crack surface roughness of X60 and UFAF Steel
R (lm)
S (lm)
H ()
X60 UFAF
8.2 24.4
39.7 63.8
17.9 40.6
The measuring range of the profilomater is 50 lm and the resolution is 0.1 lm. The mean height of profile irregularities (R), the mean spacing of the profile irregularities (S) and the mean slope of the profile (H) were statistically measured. In every sample, five sampling locations were measured with the traverse length of the stylus of 2 mm and the traverse speed of 1 mm/s in each sampling location. The final results were averaged from these five sampling locations, as given in Table 4. 3. Results 3.1. Effect of a high angle grain boundary (HAGB) on crack propagation In Fig. 3(a), there is a high-angle grain boundary (HAGB) between grain A and grain B. The loading axis is along the vertical direction. A crack propagated in grain A from right to left in a zigzag fashion. As the sample was
Fig. 3. Crack propagating through a high angle grain boundary (HAGB): (a) a crack propagated from grain A to grain B; (b) the dislocations in grain B moved under loading; (c) the crack passed through HAGB and changed its propagating direction; (d) the crack propagated in grain A along the slip direction of dislocation, and a clear dislocation free zone (DFZ) formed in front of the crack tip.
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strained, dislocations in grain A moved and piled up at front of the HAGB. When loading was increased, the dislocations in grain B started to slip (Fig. 3(b)). It can be observed that the slip direction in grain B was different from that in grain A, and the misorientation between A and B was about 30. With further increase in loading, the crack propagated across the HAGB with a significant deflection of its trajectory. In addition, a dislocation free zone (DFZ) was formed at the tip of the crack (Fig. 3(c)). After propagating in grain B for some distance, the propagation direction was gradually changed to follow the direction of the DFZ (Fig. 3(d)). 3.2. Effect of low-angle grain boundary (LAGB) on crack propagation Fig. 4 showed the process of crack propagation across a low-angle grain boundary (LAGB). In this case, grains A and B were separated by a LAGB. The loading axis was along the vertical direction. Under loading, a crack propagated from the upper right corner towards the lower left in grain A (Fig. 4(a)). As the loading increased, dislocations were generated at the LAGB in grain B, and slipped towards the left side. It was observed that the dislocation density in this grain was far higher at the higher strain
(Fig. 4(b)) than that at the lower strain (Fig. 4(a)). When the loading increased further, the crack propagated close to the LAGB, and the dislocations kept being generated from the LAGB. Some dislocations that piled up at the LAGB in grain A seemed to slip over the grain boundary to grain B. It would be noticed that the crack propagation in grain A and dislocation glide in grain B were going in almost the same direction (Fig. 4(c)). Finally, the crack propagated across the LAGB suddenly, and the propagation direction of crack in grain A and grain B did not change perceptibly (Fig. 4(d)). 3.3. Effect of LAGB with an M/A film structure on crack propagation Fig. 5 shows a crack propagating across a special LAGB. The loading axis was still along the vertical direction. Grains A and B were separated by a LAGB. It should be noticed that the morphological character of the LAGB is quite different from that in Fig. 4. The thickness of this LAGB was larger than that of a normal LAGB, and some fine structure can be observed (Fig. 5(a)). As per the previous report of the authors on microstructure characteristics of acicular ferrite [11], this kind of structure was proved to be a continuous film with thickness of several to tens of nanometers at the boundary between
Fig. 4. Crack propagating across low angle grain boundary: (a) low angle grain boundary between grains A and B; (b) crack propagating and boundary emitting dislocation; (c) crack approaching to the boundary; (d) crack propagating across the boundary.
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Fig. 5. Crack propagating across the low angle grain boundary with M/A film structure: (a) low angle grain boundary between grains A and B; (b) crack approach to boundary; (c) crack deflecting at the boundary; (d) crack propagating across the boundary.
acicular ferrite grains (Figs. 6(a) and (b)). For determining the structure of the film, TEM micro-diffraction and X-ray diffraction (XRD) were adopted. The results showed that the thin film is mainly a body-centered cubic (bcc) structure, which in steel is well known to be either ferrite or martensite (Fig. 6(c)). Therefore, this thin film should be composed of martensite and retained austenite (M/A), which transformed from retained austenite during the simulated coiling process after accelerated cooling. This M/A film in the UFAF pipeline steels, somewhat similar to the thin film reported in other conventional microstructures such as the welded AF microstructure or bainite [12,13], is significantly different in morphology from the second phase in the AF conventional carbon pipeline steels, in which the M/A component is distributed as islands. Since no visual result of the effect of M/A film structure on defor-
mation and crack propagation was reported before, it would be of great interest to investigate the process of crack propagation across such a structure. In Fig. 5(a), a crack was seen to propagate in grain A from the right side to left under loading. With increased loading, the crack propagated close to the LAGB, but no significant change occurred in the dislocation distribution in either grain A or grain B (Fig. 5(b)). When loading was further increased, the crack propagated across the LAGB, and the propagation direction of the crack deflected markedly upon crossing the LAGB (Fig. 5(c)). After the crack had propagated some distance in grain B, it was obvious that the misorientation of crack propagation between grains A and B was close to 90 (Fig. 5(d)). According to crystallographic analysis of dislocation movement and crack propagation, the slip direction in
Fig. 6. Film structure at the grain boundary of acicular ferrite: (a) bright field of M/A film structure at the grain boundary; (b) dark field of M/A film structure at the grain boundary; (c) micro-diffraction pattern of M/A film.
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bcc crystals is Æ1 1 1æ, which has fourfold symmetry [14]. Therefore, the misorientation of the slip direction up to 90 may be just the different slip direction of the same slip system. 4. Discussion 4.1. Mechanism of RICC Our previous work showed that the FCGR of UFAF steel in the Paris regime [15] was significantly less than that of X60, and also a little bit lower than that of the steel with an AF microstructure [2]. For example, at DK = 35 MPa m0.5, the FCGR of UFAF was almost half of that of X60, as shown in Fig. 7. To investigate the mechanism for the resistance against fatigue crack propagation of UFAF, trajectories of fatigue crack propagation of various samples were observed, as shown in Fig. 8. It can be seen that the crack in the X60 sample did not show much deviation and remained quite close to a straight crack. On the other hand, the crack in
da/dN, mm/cycle
X60 AF UFAF
1E-3
1E-4
30
40
∆K, MPa.m
50
60
70
80
90
the UFAF steel kinked and meandered significantly, which could lead to a rougher crack surface with more asperities. The deflection of the crack in the AF sample was intermediate between that in the X60 and UFAF samples. Therefore, it is further suggested there are some special mechanisms in the UFAF that can affect the crack propagation behavior by deflecting the crack. According to the results of the in situ TEM observations, the effect of a HAGB on deformation and crack propagation was strongly different from that of a LAGB. The change in operating slip system, which is necessary as a crack crosses a grain boundary, results in the grain boundaries being barriers for dislocation motion. For a HAGB, the formation of a DFZ in a neighboring grain will depend on the operation of its own slip system, which will lead to the significant difference of direction and crack length between the two sides of the grain boundary. Thus the crack propagation direction will change obviously after crossing the grain boundary, and the larger the misorientation between two grains, the more pronounced will be the deceleration in crack-growth rate at the grain boundary [16]. Conversely, for an LAGB, the small difference in orientation between neighboring grains would produce a relatively small retardation. Brown [17] also claimed that an LAGB fails to provide a significant barrier to crack propagation. However, the observations from in situ TEM indicated that in the UFAF steel, the LAGB with an M/A film had a similar effect to the HAGB on the dislocation motion and crack propagation. It indicated that even at an LAGB, the M/A film can be an effective barrier to dislocation motion and crack propagation, despite the small misorientation between neighboring grains. Consequently, in the UFAF it can be concluded that not only the crystallographic misorientation between two neighboring grains, but also the character of the grain boundary itself affected the crack propagation. The M/A film at grain boundary can increase the possibility of crack deflection, which leads to larger asperities than would occur in a PF+P structure as in X60. 4.2. Estimating the effect of RICC by mathematical models
0.5
Fig. 7. Fatigue crack propagation rates of X60, AF and UFAF.
As described in the above section, the UFAF pipeline steel exhibited better fatigue properties than coarse PF+P
Fig. 8. Trajectories of fatigue crack propagation of various samples: (a) coarse polygonal ferrite plus pearlite; (b) acicular ferrite; (c) ultra-fine acicular ferrite.
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Fig. 9. Scheme of the crack closure mechanism.
and coarse AF, and this improvement was the result of increased roughness in the crack surface. It is of interest, therefore, to analyze the resistance to fatigue crack propagation of a deflected crack in terms of present mathematical models and hypotheses based on the RICC mechanism. A quantitative assessment of the RICC level is possible by using the statistical modal based on a simple structural analysis [18]. The two-dimensional zigzag crack paths characterized by an average angle h, which is considered to be the determining factor for the direct effect of crack deflection, is assumed to produce the RICC according to the scheme in Fig. 9. During the loading part of the cycle, the crack tip is opened in mixed mode reaching the maximum crack tip opening displacement denoted as dmax at the peak stress (Fig. 9(b)). This displacement can be considered to be composed of both the reversible normal component d1 and the irreversible shear component d2 defining the local mixed mode (I + II). During the unloading phase, consequently, the shear displacement is not recovered so that the crack surfaces come into the contact with each other before the applied stress becomes zero [19], as shown in Fig. 9(c). The crack path deflection reduces the crack growth rate in two ways: (i) it causes a direct reduction of the local driving force for crack propagation and an increase in the total length of crack path [7]; and (ii) it causes a reduction of DK experienced at the crack tip by RICC from the nominally applied DK to the effective DKeff. These two effects are referred to as the extrinsic retardation mechanisms [20]. For the fatigue crack growth at a loading ratio of 0, the
relationship between DK and DKeff can be demonstrated as [13,21]: sffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi) ( v tan h 2 h DK; ð1Þ DK eff ¼ cos 1 2 1 þ v tan h lI 3 Rs 1 ; ¼ lII 2 Rs þ 1 1 ; Rs ¼ cos h
v¼
ð2Þ ð3Þ
where v and Rs are dimensionless parameters representing the mismatch between the crack surfaces and the roughness of the fracture surface, respectively. According to the data of crack surface roughness, the mean slope of crack profile (H), i.e., the average deflection angles h of X60 and UFAF are 17.9 and 40.6 (Table 4), respectively. Substituting the h values into Eq. (1), the effective crack propagation driving force for two microstructures would be: DK effðX60Þ ¼ 0:87DK; DK effðUFAFÞ ¼ 0:54DK. The results are consistent with the experimental data, which indicates that, due to RICC caused by crack deflection during cyclic loading, the DKeff of UFAF is much lower than that of X60. Substituting DKeff for DK in Fig. 7, the normalized fatigue crack growth rates of X60 and UFAF show that the fatigue growth rate of X60 would be lower than that of UFAF if the effect of RICC was eliminated. That is, X60 possesses a high intrinsic resistance
Fig. 10. Trajectories of fatigue crack of (a) X60 and (b) UFAF.
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∆κ eff-da/dN, UFAF ∆κ eff-da/dN, X60
da/dN, mm/cycle
∆κ -da/dN, X60 ∆κ -da/dN, UFAF
1E-3
1E-4 20
30
40
∆K, MPa.m
50
60 70 80 90
0.5
Fig. 11. Fatigue crack growth rate curves of X60 and UFAF vs. DK and DKeff.
against fatigue crack propagation resulting from matrix properties (Fig. 11). Therefore, in spite of the low intrinsic crack growth rates, the high observed crack growth rate can be due to a low level of crack closure. It was observed that, for the X60 sample, the deviation of the crack path from the macroscopic crack plane was about 10 lm, which is close to the PF grain size, while that of the UFAF sample was about 30 lm, which is much larger than the UFAF grain size (Fig. 10). So, it is further confirmed that RICC played a significant role in increasing the resistance against crack growth in the UFAF steel. Based on the analysis above, it is implied that the special microstructure constituent, i.e., the M/A film structure at the grain boundary can have a strong influence on deflecting the crack trajectories and enhancing RICC. Therefore, it is hoped that, in the future, more intensive study will be carried out to further develop the necessary understanding of the nature and role of the grain boundary films, in order to further improve the mechanical behavior of these high strength pipeline steels. 5. Conclusions The micro-fracture behavior of a high strength pipeline steel with a UFAF microstructure was directly observed by in situ TEM tensile testing in order to study the effect of the grain boundary structures on the FCGR. According to the results, the following conclusions can be made.
(1) A LAGB did not influence the crack propagation significantly, but an LAGB with an M/A film structure was an effective barrier to dislocation motion and crack propagation, despite the small misorientation between neighboring grains. Consequently, such boundaries can lead to strongly deflected fatigue crack trajectories, which, in turn, lead to rougher asperities on the crack surfaces during cyclic loading. (2) The mathematical analysis of the fatigue mechanism indicated that despite a poor intrinsic resistance against crack growth, owing to refined grain size, the UFAF possesses a lower effective DK (DKeff) than coarse polygonal ferrite plus pearlite and coarse AF, due to a higher level of RICC that is attributed to the increasing in asperities on the crack surfaces. Therefore, by enhancing RICC, the M/A films at grain boundaries play a significant role in the improvement of the resistance against fatigue crack growth in UFAF. (3) This RICC mechanism enhanced by the M/A film structure, which is a novel discovery for the fatigue resistance mechanism in low carbon micro-alloyed pipeline steels, provides a new idea for the development of pipeline steels with both high fatigue properties and high strength. Acknowledgments The work was financially supported by funding from the National Natural Sciences Foundation (No. 50471106). The authors thank Dr. William Warke from BP Amoco for amending the paper and Dr. Xiaoqiang Zhang from our institute for the help with experiments. References [1] Hagiwara N, Meziere Y, Oguchi N, Zarea M, Champavere R. JSME Int J Ser A 1999;42(4):610. [2] Zhong Y, Shan Y, Xiao F, Yang K. Mater Lett 2005;59(14–15): 1780. [3] Suresh S. Fatigue of materials. 2nd ed. Cambridge: Cambridge University Press; 1998. p. 257. [4] Suresh S, Ritchie RO. Metall Trans A 1982;13:1627. [5] Elber W. Eng Fract Mech 1970;2:37. [6] Lerch BA, Jayaraman N, Antolovich SD. Mater Sci Eng 1984;66:151. [7] Kitagawa H, Yuuki R, Ohira T. Eng Fract Mech 1975;7:515. [8] Campbell JP, Venkateswara Rao KT, Ritchie RO. Metall Mater Trans A 1999;30:563. [9] King JE. Mater Sci Tech 1987;3:750. [10] Vinogradov A, Nagasaki S, Patlan V, Kitagawa K, Kawazoe M. Nanostruct Mater 1999;11(7):925. [11] Zhao MC, Hanamura T, Qiu H, Ke Yang. Mater Sci Eng A 2005;395(1–2):327. [12] Spanos G, Moon DW, Fonda RW, Menon ESK, Fox AG. Metall Mater Trans A 2001;32:3043. [13] Chang LC, Bhadeshia HKDH. Mater Sci Technol 1995;11:874. [14] Hull D, Bacon DJ. Introduction to dislocations. 3rd ed. Oxford: Pergamon Press; 1984. p. 257. [15] Paris PC, Tada H, Donald JK. Int J Fatigue 1999;21:S35. [16] Chen D, Sixta ME, Zhang XF, De Jonghe LC, Ritchie RO. Acta Mater 2000;48:4599.
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