2021 Volume 62 Issue 4 Pages 570-573
Temper embrittlement in high-strength steel occurs due to the segregation of impurity elements along prior austenite (γ) grain boundaries, resulting in a decrease in intergranular strength. However, the grain boundary properties have not yet been experimentally investigated. In this study, we developed a micro testing method to selectively process micro-cantilever specimens based on a specific prior γ-grain boundary to measure the grain boundary fracture toughness. Fracture toughness tests using micro-cantilever specimens with side grooves and notch were carried out at 183 K and brittle fracture behavior was successfully obtained. The fracture surface showed almost flat and brittle crack propagation along the prior γ-grain boundary. The derived fracture toughness value, KQ, is lower than that from previous values. By employing this method for the evaluation of intergranular fractures in steel, we believe that we can evaluate more intrinsic prior γ-grain boundary properties compared with conventional methods.
Steel used in offshore construction requires high strength to cope with the large structures. In general, temper embrittlement can occur in intergranular fractures along prior austenite (γ) grain boundaries in high-strength steel utilizing tempered martensite; this results in a significant loss of toughness. This type of intergranular fracture has been explained by the segregation of impurity elements, such as phosphorus (P) into the prior γ-grain boundaries during tempering, inducing a decrease in intergranular strength.1,2) In addition, the toughness value has been reported to decrease with an increase in the amount of P segregation, and the amount of decrease in grain boundary strength has been estimated from the corresponding decrease in the toughness value.3,4) However, grain boundary properties like strength and fracture toughness have not yet been experimentally investigated.
To date, the evaluation of grain boundary properties has been carried out using specimens with a notch introduced into the interface of the bicrystal. However, because it is almost impossible to reproduce prior γ-grain boundaries using bicrystals, another method is needed. One such evaluation method is micro mechanical testing, and various studies have been conducted to evaluate the strength and fracture properties of micro constituents in materials, including phase boundary fracture strength.5,6) Therefore, fracture toughness tests for micro size specimens that selectively process a specific prior γ-grain can reveal intrinsic grain boundary properties. In fact, several research groups have attempted to evaluate grain boundary properties.7–10) However, in fracture toughness tests, if the fracture toughness value (K) evaluated on different size specimens are the same, then the plastic zone at the notch tip is the same size. A smaller specimen will not satisfy the small-scale yielding condition. Therefore, brittle fracture may occur in conventionally sized specimens, while ductile fracture occurs in micro size specimens.11,12) For this reason, intrinsic fracture-toughness evaluation of prior γ-grain boundaries using micro size specimens has not been carried out.
In this study, we developed a low-temperature testing method for specimens with a high plastic constraint at the notch tip. Furthermore, we attempted to evaluate brittle fracture toughness of prior γ-grain boundaries using these methods.
The steel used in this study was a low-alloy steel; its chemical composition is shown in Table 1. To facilitate the formation of intergranular fractures, 0.1% of P was added. It was made by water quenching after heat treatment at a temperature of 1473 K for 3600 s. Post-etching optical microscopy (OM) observation confirmed a fully martensitic single-phase microstructure with coarse prior γ-grains with an average diameter of 746 µm. Charpy impact tests were carried using the 2 mm V-notch specimen in accordance with JIS Z2242.
The specimen shape was of the cantilever type. It was prepared by cutting a plate from the bulk material, which was then ground to a thickness of ∼250 µm using emery paper. Both surfaces of this thin foil were polished using colloidal silica paste. Scanning electron microscopy (SEM), combined with an electron backscatter diffraction (EBSD) was used at a scanning step size of 0.6 µm and an accelerating voltage of 20 kV to determine the crystallographic orientations of samples. EBSD analysis was performed using the TSL OIM software (v. 7.1.0). After determining the grain boundary, a micro-cantilever specimen 1250 µm in length, 250 µm in width, and 250 µm in thickness was fabricated using electrical discharge. Following this, a notch 80 µm in depth and side grooves 20 µm in depth were introduced along the grain boundary using a focused ion beam machining machine. Figure 1 shows the SEM image of micro-cantilever specimen. Figure 2 shows a crystallographic orientation map around the notch tip, determined by EBSD analysis, along with the (011) and (111) pole figures of regions A and B corresponding to the CP (close-packed) map. Regions A and B were presumed to have CP groups generated from different prior γ-grains. Therefore, it was confirmed that the notch was introduced along a prior γ-grain boundary.
Scanning electron microscope (SEM) image of the micro-cantilever specimen.
(a) Electron backscatter diffraction (EBSD) crystallographic orientation map around the notch tip, and (b) pole figures for the (011) and (111) of the martensite microstructure.
Micro bend tests were performed at room temperature (RT) and 183 K, in vacuum (2 × 10−3 Pa) using a micro testing machine developed by the authors (Fig. 3). The specimen are placed in a dedicated holder and then fixed on a precise α · β · θ stage. A piezoelectric element was used as the actuator for loading. A diamond spherical indenter with a 5 µm tip radius was attached to the end of the actuator via a force sensor with a maximum load of 5 N, and the indenter was pressed against the test specimen to apply a bending load. The loading mechanisms were fixed onto a precise X · Y · Z stage. The apparatus was contained within a chamber, and the specimen was cooled to the test temperature after vacuum. The test temperature was controlled to ±5 K using a cooling system with heat transfer by liquid nitrogen connected to a dedicated holder and cooling plate in the chamber. The temperature was measured with a thermocouple mounted near the specimen on a dedicated holder. Furthermore, micro bending test was performed at a displacement rate of 1 µm/s. The section around the notch was monitored via OM during the bend test. The crack tip opening displacement (CTOD) was measured using a video.
Schematic diagram of the testing machine.
The ductile-to-brittle transition temperature (DBTT) was 383 K and intergranular fracture surfaces were shown below DBTT. The high-magnification SEM images of intergranular fracture surfaces are shown in Fig. 4. A directionally aligned tongue pattern was observed on some intergranular fracture surfaces.
Scanning electron microscope (SEM) image of intergranular fracture surface after the Charpy impact test.
For the micro bending tests, when the notched specimen without side grooves was tested at RT, ductile crack propagation was observed. When the notched specimen with side grooves was tested at 183 K, the specimen fractured in a brittle manner. The load-CTOD curve at 183 K and deformation of a test specimen at each load are shown in Fig. 5 and Fig. 6, respectively. CTOD increased linearly with respect to load from 0 to ∼2.5 N, after which it began to deviate from the line. CTOD increased significantly at ∼3.4 N, and slip bands with an angle of 50° to the LD direction were observed on the side surface of the fixed end side of the notch. Finally, the specimen completely ruptured at ∼3.5 N.
Load-crack tip opening displacement (CTOD) curve.
Optical microscope (OM) image of the specimen side surface during loading.
Figure 7 shows a fracture surface after micro bend testing. The fracture surface is almost flat and no river pattern is observed. Meanwhile, tongue patterns with an angle of 25° to the ND direction can be observed. Figure 8 shows a crystallographic orientation map of the fracture surface determined by EBSD analysis, along with the (011) and (111) pole figures corresponding to the CP map. It was confirmed that the entire fracture surface had a CP group generated from the same prior γ-grain. In addition, the tongue pattern on the fracture surface was parallel to the habit plane. As cleavage cracks propagate along the (001) plane, it has been reported that tongue patterns form by bypassing the twin boundaries.13) In the case of intergranular fractures, it is considered that tongue patterns also form owing to an intersecting twin generated along the habit plane.
Scanning electron microscope (SEM) image of the fracture surface after micro bend testing.
(a) Electron backscatter diffraction (EBSD) crystallographic orientation map of the fracture surface after micro bend testing, and (b) pole figures for the (011) and (111) of the martensite microstructure.
Generally, river patterns are observed along cleavage fracture surfaces of martensitic steel,14) but not in this steel. Contrarily, the tongue pattern observed on the intergranular fracture surface, as shown in Fig. 4, was also observed on the micro bend fracture surface. In addition, the entire fracture surface was in the same prior γ-grain. On this basis, it is believed that the crack propagated along a prior γ-grain boundary.
When the load-CTOD curve deviates significantly from the initial proportional line, a stretch zone is generally observed at the notch tip; however, we observed no such ductile region (Fig. 7). When a bending load is applied, tensile stress is applied around then notch tip, while compressive stress is applied to the lower side of the ligament. In the martensite microstructure, slip occurs along the habit plane when the habit plane is at a 48° angle to the load direction.5) Here, the habit plane had an angle of 50° with respect to the compressive stress along the LD direction, and the plastic constraint of the compressive stress side was considered to be lower than that of the tensile stress side under the influence of the notch; this suggests that plastic deformation preferentially occurred on the lower side of the ligament. These results suggest that despite the increase in CTOD, plastic deformation at the notch tip did not occur and brittle cracking was initiated.
Finally, K was calculated. The crack initiation load could not be clearly distinguished from load-CTOD curve. Therefore, following ASTM E1304, we assumed a 5% offset load for the crack initiation load. For the calculation of the fracture toughness value for the side groove specimen, equations were proposed.15) Using those equations, the K value was found to be 19.8 MPam1/2. However, specimen shape and testing method do not comply with ASTM E1304. Therefore, it is not appropriate to use equation for small scale yielding to determine the validity of plane stress fracture toughness. Therefore, in this study, the fracture toughness value is expressed with KQ instead of KIC.
According to previous reports, the fracture toughness value evaluated by three point bending tests in accordance with BS7448 in tempered martensitic steel is 30.8 MPam1/2.16) Compared with the previous value, the value calculated here is lower. Whereas the previous study evaluated the entire fracture toughness values of multiple prior γ-grain boundaries, this study evaluated the fracture toughness of one prior γ-grain boundary. Therefore, this value is considered to represent the intrinsic grain boundary property.
To evaluate grain boundary fracture toughness directly, we developed an evaluation method using a micro-cantilever specimen to selectively process a specific prior γ-grain boundary. When a specimen with side grooves and a notch was tested at 183 K, we were able to simulate brittle crack propagation along the prior γ-grain boundary and obtain grain boundary fracture toughness. The fracture toughness value obtained in this study was lower than that previously evaluated using a conventional specimen size. By using this method, we believe that we can evaluate more intrinsic grain boundary properties, including the effect of the segregation of impurity elements such as P on grain boundary properties.
