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Effects of annealing on microstructure and mechanical properties of nano-grained Ni-based alloy produced by severe cold rolling
Author’s Accepted Manuscript Effects of annealing on microstructure and mechanical properties of nano-grained Ni-based alloy produced by severe cold rolling Yanle Sun, Songqian Xu, Aidang Shan www.elsevier.com/locate/msea
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S0921-5093(15)30098-8 http://dx.doi.org/10.1016/j.msea.2015.06.043 MSA32481
To appear in: Materials Science & Engineering A Received date: 26 April 2015 Revised date: 12 June 2015 Accepted date: 13 June 2015 Cite this article as: Yanle Sun, Songqian Xu and Aidang Shan, Effects of annealing on microstructure and mechanical properties of nano-grained Ni-based alloy produced by severe cold rolling, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2015.06.043 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Effects of Annealing on Microstructure and Mechanical Properties of Nano-grained Ni-based Alloy Produced by Severe Cold Rolling
Yanle Sun1, Songqian Xu2, Aidang Shan1*
1. School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai, 200240, People’s Republic of China 2. Baosteel Special Steel Co., Ltd. Shanghai, 200940, People’ Republic of China
Corresponding author
Tel.: +86 21 54747556; fax: +86 21 54740825 E-mail: [email protected] (Aidang Shan)
1
Abstract Nano-grained (NG) Ni-based alloy with average grain size about 50 nm was produced via severe cold rolling at room temperature. Compared to the solution-treated alloy prepared for rolling, the cold-rolled (CRed) NG alloy exhibited approximately 475 %, 130 % and 180 % increase in the average yield strength (YS), ultimate tensile strength (UTS) and microhardness, respectively. Effect of the annealing on mechanical properties and microstructure of this NG Ni-based alloy were investigated systematically. The annealed-CRed alloys possessed NG (90 nm) and ultrafine-grained structure (200 nm) after 60 min annealing at 700 and 800 oC, respectively, suggesting high thermal stability of the nanostructure. Peak YS and UTS of 2047.5 and 2142 MPa, respectively, were obtained with appropriate annealing temperature and time. The mechanisms of variations in mechanical properties were analyzed based on phase composition and microstructure evolution of the annealed NG Ni-based alloy. Both the high thermal stability and strength are due to the formation of the γ′ precipitates and slight grain growth of the NG matrix.
Key words: Nano-grained (NG); Ni-based alloy; Severe cold-rolling; Mechanical properties; Microstructure; Thermal stability
2
1. Introduction Bulk nano-grained (NG) metals, having a grain size in the nanoscale (10-100 nm), attract considerable attention due to their exceptional physical, chemical and mechanical properties relative to their coarse-grained counterparts [1-4]. Severe plastic deformation (SPD) offers an attractive route to fabricate bulk NG materials. The microstructure and mechanical properties of NG and UFG Ni-based alloy prepared by different SPD processes, such as high-pressure torsion (HPT) [5-7], multiple forging (MF) [5-7] and other special SPD processing [8-10], have been investigated by some researchers. In comparison with the ECAP and HPT, severe cold-rolling process could overcome the major limitations of low productivity of the former and small work-piece size of the latter [11, 12]. While there were some research on the NG pure nickel [13, 14] and single phase Ni-based alloy [15] obtained by severe cold-rolling, there were no study on the microstructure and mechanical properties of a dual-phase NG Ni-based alloy prepared by cold-rolling. Hence, there is a strong scientific incentive to systematically study the evolution of microstructure, phase constitutions and variation of mechanical properties of cold-rolled (CRed) dual-phase NG Ni-based alloy during annealing. The Ni-based alloy in present research consists of a solid solution γ matrix strengthened by stable nanoscaled spherical γ′ phase (Ni3 (Al,Ti)), ordered face centered cubic L12 structure [16]. Our research is motivated by two objectives: first, to use severe cold rolling process to fabricate an NG Ni-based alloy; and second, to systematically investigate the effect of annealing on mechanical properties, phase composition and microstructure of this NG Nibased alloy. The grain size, crystal preferred orientation, phase composition, microhardness as well as tensile properties of the NG alloy were characterized after 3
different annealing process and a high-strength, thermally stable dual-phase nanostructured Ni-based alloy through precipitation were obtained. 2. Experimental procedures The starting material was bulk microcrystalline (20~30 µm, seen in Fig. 1) Ni-based alloy with dimensions of 18 × 15× 100 mm3 and the alloy compositions were given in Table 1. Thin sheet samples of NG alloy with a thickness reduction of 98 % from 18 mm to 0.3 mm was produced by severe cold rolling at room temperature. Prior to this, the bulk materials prepared for cold rolling were solution treated (STed) at 1100 oC for 2 h to ensure that all the initial precipitations were dissolved and a homogeneous microstructure was created. Phase constitutions of the STed, CRed and annealed alloys were determined with Xray diffraction (XRD) technique, using Cu K (𝜆�� = 0.1540562 nm) radiation on a Shimadzu XRD-6000 diffraction instrument, operating at 40 kV and 30 mA between 20 and 100 deg (2θ) at a step of 0.02 ° and a counting time of 0.6 s per step. Transmission electron microscope (TEM) and selected area electron diffraction (SAED) studies were performed using a JEM-2100F operating at 200 kV. The observed sections of CRed and annealed samples for TEM were ND-TD plane. Foils were cut, punched to 3-mm-diameter discs, mechanically polished to a thickness of 50 μm and finally twin-jet polished using an electrolyte of 5 % perchloric acid and 95 % alcohol that was cooled to 243 K with an applied voltage of 40 V. RD, TD and ND stand for rolling, transverse and normal directions, respectively. Specimens prepared for microhardness tests were machined from the CRed-sheets and subsequently annealed at temperatures from 500 to 900 oC for different length of 4
time. Microhardness test was carried out on the ND-RD plane with a load of 500 g and a dwell time of 15 s using the Buehler Micromet Hardness Tester with a Vickers diamond pyramid indenter. For each measured plane, at least 9 points were measured to calculate the mean value of microhardness and the measurement error was no more than 2.5 %. Flat specimens for tensile testing, with a thickness of 0.3 mm, were sectioned along the longitudinal rolling direction from the CRed-sheets by electrical discharge machining. Thickness of tensile specimens cut from the STed bulk material which was prepared for cold rolling was 1 mm. The gage length and width of all the tensile specimens were 15 and 3 mm, respectively. All the tensile tests were conducted at ambient temperature by using a Shimadzu AG-10kNA tension tester at a strain rate of 10-3 s-1. 3. Results 3.1 Phase components and microstructure 3.1.1 XRD analyses Fig. 2 shows the XRD patterns of the STed Ni-based alloy prepared for cold rolling as well as the as-CRed state. All the diffraction peaks corresponded well with those of a pure nickel standard (JCPDS card 01-087-0712), which confirms that the STed Ni-based alloy solely consisted of γ phase and no phase transformation occurred during subsequent cold-rolling deformation. The intensity ratios of all the diffraction peaks in the STed alloy were well in accordance with those of the standard. However, the intensity ratios of (2 2 0) peak of the CRed alloy was much larger than those of STed sample, suggesting preferred orientations along the (2 2 0) crystal plane in the CRed material. Additionally, a slight peak shifting from the pure nickel position to smaller Bragg angles was present in the XRD patterns [17], which was consistent with the fact that both the Ni-based alloys 5
were in solid-solution state. Interestingly, the preferred orientation along the (2 2 0) crystal plane slightly increased with the increase of annealing temperature as illustrated in Fig. 3. From Fig. 3 (b)-(e), new diffraction peaks at 2 of 35.5°, 63.7° and 80.4° were present, which were corresponding to the (1 1 0), (2 1 1) and (2 2 1) crystal planes of γ′ phase. It can be ascertained that the second phase of γ′ precipitated after annealing for 1 h at 500 , 600, 700, 800 oC, respectively. The intensity of (1 1 0), (2 1 1) and (2 2 1) peaks of γ′ phase became stronger with the annealing temperature increasing from 500 to 700 oC and then reduced with the temperature further increasing to 800 oC. The variation of intensity of diffraction peaks of γ′ phase showed that the volume fracture of γ′ precipitation increased gradually as the annealing temperature increasing from 500 to 700 oC and then significantly decreased as the annealing temperature further increasing to 800 oC. 3.1.2 TEM analyses TEM analyses prove that cold rolling process of the Ni-based alloy resulted in complete refinement of the initial coarse-grained structures into nanoscale ones. It is evidenced in Fig. 4 by the bright- field (BF) and dark-field (DF) images as well as the diffraction rings. The TEM images clearly show that the most common feature in the asCRed sample is a random distribution of grains with average size approximately of 50 nm, as measured in Fig. 4 c and d. The microstructure is composed primarily of equiaxed grains, along with some distorted structures. Inspection of BF and DF images shows the presence of high densities of dislocations and that the grain boundaries tend to be curved, irregular and/or ill-defined, which reveals that the grain boundaries introduced by the severe cold-rolling were severely distorted and in a high-energy non-equilibrium 6
configuration. In addition, the SAED pattern in Fig. 4 b exhibits typically continuous rings ascribed to the NG γ phase, which is electron diffraction characteristic of ultrafine grains separated by boundaries having high-angle of misorientation due to the transformation from dislocation cell and sub-grain boundaries as increasing strain induced into the sample [18]. All of these characteristics are typical of materials produced by SPD [19-21]. The CRed samples after annealing for 1 h at 700 and 800 oC, were respectively examined in the TEM to track the evolution of the nano-grains produced by severe cold rolling. Fig. 5 presents the TEM images of CRed samples after annealing at 700 oC for 1 h. The average grain size of the γ matrix is nearly 90 nm, as measured in Fig. 5 b, suggesting a slight increase of the matrix grains in compared with that of the as-CRed material. It is evident that most grains are equiaxed shapes, but with sharp grain boundaries and the morphology of the grains is similar to that of the as-CRed sample. However, grain boundaries in the samples after such an annealing become more distinct, which ascribes to the decrease in the number of defects in near-boundary regions and some extent of stress relaxation. SAED pattern clearly shows the coexistence of� γ and γ′ phases, which is consistent with the XRD analysis in “XRD analyses” section. Additionally, vast scale dislocations can be observed in the matrix and the SAED pattern remains to be continuous rings ascribed to the nano-grains. Fig. 6 presents the TEM images of CRed samples after annealing at 800 oC for 1 h. The average grain size of γ matrix significantly increases beyond the regime of nanoscale and is around 200 nm, as measured in Fig. 6 b. Due to the grain growth, the SAED pattern starts to present separate spots, which further confirms that the grains have grown 7
up apparently. Nevertheless, the grain size distribution emerges a bimodal distribution, as there are several grains below 100 nm. Fig. 6 a and b show that the morphology of γ grains are in equiaxed shape with distinct grain boundaries, which is a structure characteristic with lower density of defects on the grain boundary in comparison with the as-CRed sample. In the structure, the present of thin straight grain boundaries with a junction at an angle of 120° exhibits the occurrence of recrystallization. The fringe contrast of several grain boundaries in both BF and DF images also reveals the intense occurrence of recovery in these grains and that their state verges to the equilibrium. The SAED image in Fig. 6 c clearly shows the existence of γ′ phase. 3.2 Mechanical behavior 3.2.1 Microhardness Fig. 7 presents the variation of microhardness of annealed-CRed samples as a function of the annealing temperature while the annealing time was fixed to be 1 h. It can be seen the microhardness of the as-CRed sample was about 480 Hv (kgf·mm-2), which was almost three times of the solution-treated microcrystalline Ni-based alloy before rolling, of which microhardness was about 169 Hv. As shown in Fig. 7, microhardness of samples increases firstly and then decreases as annealing temperature rising from 500 to 900 oC and a peak microhardness up to 610 Hv was attained at 700 oC. When the annealing temperature increased to 800 oC, the microhardness was still slightly higher than that of the as-CRed alloy. However, an abrupt drop of the microhardness from 480 to 330 Hv suggested a rapid softening of the CRed material when annealing at 900 oC. Fig. 8 shows that the gradually decreased of microhardness as the annealing time prolonged from 1 h to 100 h at 800 oC. 8
The tendency of the microhardness that firstly increased sharply and then fell slightly both at 600 and 700 oC with the annealing time. As shown in Fig.9, rapid increase in microhardness was observed during the subsequent annealing. Even annealed for 5 minutes, a dramatic increase in microhardness was resulted and the incremental value of microhardness was much higher at 700 than 600 oC, while approximately increment higher at 600 than 500 oC. The peak microhardness over 610 Hv was achieved when annealed for 1 h at 700 oC, while for 3 h at 600 oC and 10 h at 500 oC. The annealed-CRed Ni-based alloy exhibits an outstanding microhardness in the γ´strengthened Ni-based alloys. The value of microhardness remained to be essential constant above 600 Hv onset 10 minutes till 2 h and hold 570 Hv as annealing time prolonged to 10 h at 700 oC. As the annealing treatment extended to 300 h at 700 oC in Fig. 10, the microhardness was observed to be gradual decrease with the annealing time. Microhardness of 440 Hv was resulted and this value of microhardness was slightly lower than that of the CRed alloy but 38 % greater than that of the peak-aged (320 Hv) alloy in the conventional microcrystalline state. Furthermore, the top value of microhardenss over 600 Hv hold for a longer duration at 600 oC, starting from 2 to 10 h. The high and stable microhardness of the annealed alloys demonstrates high thermal stability under 700 oC. 3.2.2 Tensile behavior The representative engineering stress-strain curves established from samples machined from STed and CRed Ni-based alloys before and after various annealing processing are presented in Fig. 11. The average values of more detailed mechanical properties are summarized in Table 2. In the STed state, the average yield strength (YS), 9
ultimate tensile strength (UTS) and elongation of coarse-grained samples are 253 MPa, 684 MPa and 44 %, respectively. When the STed alloys annealed at 700 ℃ for 1 h, the average YS and UTS are 79.6 % and 31 % higher than those corresponding to the STedstate, respectively, and the elongation is slightly lower (454.5 MPa, 898 MPa and 41.1 %, respectively). In the case of the CRed state, it is evident that the strength was greatly enhanced after cold-rolling processing. The YS and UTS of the NG Ni-based alloy in the CRed state were 1445 MPa and 1557 MPa, respectively, which were respective over five times and two times of the coarse-grained Ni-based alloy in the STed state prepared for cold-rolling process. However, as predicted, the ductility of the alloy decreased largely after cold-rolling [3, 22]. Strengthening through severe cold-rolling process results in a material with a tensile curve that peaks immediately after yielding (curve of as-CRed). Such a trend of strengthening accompanied by a loss of ductility is generally true for various materials processed by SPD ways [14, 21, 22]. Based on the average value of mechanical properties for all samples machined from a processd alloy, the YS and UTS further increased in the annealed-CRed Ni-based alloys. When the CRed alloys were annealed at 600 oC for 1 h, as shown in Fig. 11 and Table. 2, the YS and UTS were 2013 and 2061 MPa, respectively. What is more, a higher YS and UTS up to 2047.5 and 2142 MPa were produced as the annealing time prolonged to 5 h. However, the high strength was accompanied by ductility minimum. The high YS and UTS, together with the microhardness of the annealed-CRed alloys suggested stable mechanical behavior after annealing at 600 oC. When the annealing temperature was fixed at 700 oC, rapid enhancement of the YS and UTS by 33 % and 28.9 % were resulted, respectively, as annealing for 10 min (1939 and 2007.5 MPa, respectively). The 10
YS and UTS were 1837 and 1927 MPa, respectively, as annealing for 5 h. When the annealing time extending from 10 min to 5 h, the YS and UTS are holding above 1800 and 1900 MPa, respectively. However, the elongation is limited and slightly increased with the increase of annealing time. When the factors of the annealing temperature and time was fixed, alternatively, and the other one was considered as variable factor, the nature in change of YS and UTS was quite in accordance with that of the microhardness. 4. Discussion 4.1Microstructure and Mechanical properties of the CRed Ni-based alloy. 4.1 Strengthening mechanism of CRed Ni-based alloy. In the present work, the bulk NG Ni-based alloy with average grain size of 50 nm was fabricated by severe cold-rolling processing. X-ray diffraction (XRD) analysis of asCRed sample suggested that all the additional alloy elements were completely dissolved into the Ni matrix, as only diffraction peaks of the γ phase were detected. Thus the possible strengthening effect of solute atom clusters or precipitates can be neglected. Increase of the microhardness and strength in the as-CRed alloy mainly results from the grain boundary strengthening. High fraction grain boundaries surrounding the nanograins act as obstacles to the movement of the dislocations which would pile up near the grain boundaries, leading to a steep enhancement in the microhardness, YS and UTS [11]. TEM observation revealed that the NG matrix which typically composed of grains around 50 nm were produced by the severe cold-rolling processing. Moreover, the high defect densities inherent in the deformed materials make a further contribution to the strength. Hugnes and Hansen had testified that cold- rolling process of nickel to a strain of 90 % 11
resulted in a greater increase of strength than that calculated in the Hall-Petch rule [13, 23]. Similar results were verified in other materials such as Al alloys by high pressure torsion (HPT) [24]. Since the initial establishment of the Hall-Petch equation was based on the un-deformed materials, it was reasonable that the low angle grain boundaries (LABs) and non-equilibrium state of grain boundaries (NGBs) were not counted, which were only abounded in the severe plastic deformed materials. But in the materials processed by SPD techniques, the NGBs and LABs introduced into the material during processing would play an considerable contribution to the strength [25]. In case of the CRed Ni-based alloy in our research, the contribution of these two boundaries are taken into account and the total strength of the CRed samples can be estimated as followed [26, 27] :
σ y ~ 0 kd -1/2 LAB NGB
(1)
Where 0 is the threshold stress, k is the Hall-Petch constant, d is grain size, kd -1/2 � represents the strengthening contribution of the high angle grain boundaries (HAB) between grains. The LAB and NGB represent the strengthening contribution of the LABs and NGBs, respectively. TEM observation in Fig. 4 reveals that the microstructure is basically composed of HABs, which make a major contribution to the strengthening in the as-CRed samples. From the Eq. (1), it is evident that the ultra-high YS and UTS of the as-CRed Ni-alloy are mainly resulted from the extremely fine grains and additional enhanced by high defect densities [28]. In the case of the microhardness, the Hall-Petch relationship in Eq. (1) can be applied and its values significantly increased from 169 Hv of the coarse-grained STed alloy to 480 Hv of the NG CRed-alloy. 12
4.2 Improved strength of the NG Ni-based alloy via annealing When the as-CRed Ni-base samples annealed at 500, 600 and 700 oC, further increase of the YS, UTS and microhardness were resulted. No doubt the very rapid hardening is mainly due to the rapid formation of the nanoscaled γ′ precipitates in the NG matrix while grain growth of the NG matrix is slight at these temperatures. This is confirmed by the TEM observations in Fig. 5, “XRD analysis” in Fig. 3 and the microhardness variation in Fig. 9. Higher microhardness and strength correlate with higher volume fraction of the γ′ precipitates. While we have not explicitly imaged the evolution of the nanoscaled precipitates in the alloy, such rapid hardening in nanostructured precipitation-strengthening alloy is known to be due to the rapid formation of the second-phase precipitates [9, 29]. Microhardness of the CRed alloy increased from 480 to 615 Hv while the corresponding YS and UTS increased from 1455 and 1557 to 2047.5 and 2142 MPa, respectively. Generally, the correlation between the YS and microhardness has been determined to be [30, 31]:
Hv c y
(2)
Where y is the uniaxial YS and c is the elastic constraint factor. For blunt pyramid indenters, it was reported that c 0.3 when the hardness and yield strength are measured in kgf·mm-2 and MPa [30, 31] . According to calculation of microhardness values on the ND-RD plane of the as-CRed and annealed-CRed samples, the factor of c was slightly waved for the different states of material. Since both the microhardness and mechanical strength are significantly affected by microstructure, the correlation between the YS and microhardness are also affected by microstructure [32]. In the as-CRed alloy, the high 13
density of defects such as dislocations were expected to affect this correlation and led to a large value of factor to 0.33. When the as-CRed alloy annealed in various conditions, a slight fluctuation of the value of factor c was resulted for the evolution of microstructure during annealing. According to the SAED, TEM and XRD analysis, such annealed-CRed alloys were composed of the NG matrix of γ phase, typically embedded with dispersed γ′ precipitation. In this precipitate-strengthened NG Ni-based alloy, strengthening is due to the stable ultra-fine γ grains and the dispersed of nanoscaled γ′ precipitates. Thus, the effective strengthening mechanism are the Hall-Petch strengthening due to the NG matrix and the Orowan mechanism due to the spherical γ′ precipitates. Thus, resulting strengthen (τ y ) in the annealed-CRed Ni-based alloy can be approximately expressed as followed [9]:
τy y
Gb
λ
(3)
Where y �is the strength of the as-CRed alloy showed in Eq. (1),� G is the shear modulus, b �is the Burgers vector and λ �is the mean spacing between γ′ particles. The part of��
Gb
λ
in Eq. (3) represents the effect of Orowan strengthening attributed to the γ′
particles, which suggest that dislocations circumvent hard particles leaving dislocation loops behind. Since d 1/2 �(in� Eq. (2))� and� λ �(in� Eq. (3))� are in an order of magnitude, these two strengthening effects are comparable and either one can be neglected. Thus, the resulting strength is the tradeoff between the effect of precipitation strengthening and matrix softening. When the CRed alloy annealed at 800 oC for 1 h, recrystallization of the 14
structure occurred as the present of distinctly straight grain boundaries with a junction at an angle of 120° and the grains significantly grown to about 200 nm. The softening effect of the recrystallization and grain growth prevailed over the strengthening effect of the precipitation as the annealing time prolonged over 2 h at this temperature. Thus the strength was observed to be decrease. From the Orowan formula (
Gb
λ
), it can be seen that the finer the precipitate size
and the larger their volume fraction, the greater effect of precipitation strengthening produced. In the “XRD analysis” section, it is clear that the intensities of diffraction peaks of γ′ phase become much stronger as annealing temperature increasing from 500 to 700 oC, showing evident increase in the volume fraction of γ′ precipitation. Foreparts of curves (forward the value of 60 on the horizontal axis) in Fig. 9 suggested that increase of annealing temperature from 500 to 700 oC would accelerate the rapid hardening of the γ′ precipitates. Evolution of microhardness of annealed-CRed samples was in accordance with the XRD results. When the CRed alloy annealed under 800 oC, strength softening of the NG matrix were more apparent for the growth of grains and annihilation of a part of LABs and NGBs and the strengthening effect of nanoscaled γ′ precipitates would go down for less volume fraction of the precipitation. It is reported that a peak hardness value of 630 Hv was achieved in the γ″ and γ′strengthened NG� Inconel 718 and the microhardness of the counterpart with coarse grains after solution treatment was 300 Hv [9]. In the present work, a peak value of microhardness of this γ′ precipitate-strengthened NG Ni-based alloy was 615 Hv, almost by a factor of four higher than that of the coarse-grained alloy in the STed state, which 15
was about 169 Hv. Comparison of the above results, it can be seen much more significant enhancement of the microhardness was achieved in this Ni-based alloy. 4.3 High thermal stability In the present work, a notable thermal stability of the NG CRed-alloy was observed. Most recently study shows that the critical temperature of nanostructured Inconel 718 with grain size of 100 nm produced by plane strain machining is 0.6 Tm (Tm, melting point in K) [9]. It has been reported that NG Ni-20%Cr and Inconel 718 produced by HPT, typically with a mean grain size of 50-80 nm, are thermally stable up to 0.46 and 0.5 Tm, respectively [7]. The temperature threshold for the retention of nanocrystalline structure of the Inconel 718 is higher than that of the Ni-20%Cr, which is due to the effect of the uniformly distributed coherent disk-type-γ′′-precipitates within grains in the Inconel 718. Improvement of thermal stability of the nanocrystalline alloy by introducing a dense dispersion of fine second-phase precipitates has also been reported in many other alloys such as Ti alloys and Al alloys [29, 33-35]. In the present research, the Tm� of the Ni-based alloy is 1300 oC (1573K) that calculated by the JMatPro software. TEM observation of the CRed and annealed-CRed alloys shows that the average grain size maintains below 100 nm when annealed for 60 min at 700 oC and thus this NG matrix with grain size about 50 nm is stable up to 0.62 Tm. In the view of above discussion, the high thermal stability is closely associated with the pinning effect of the� γ′ precipitates on grain boundaries and thus grain growth can be significantly blocked. Further research on the thermal stability of NG Ni-based alloy is undertaken. 5. Conclusion 1. NG Ni-based alloy with average grain size of about 50 nm was prepared via 16
severe cold-rolling at room temperature. As compared to the STed alloy before rolling, the CRed alloy exhibits approximately 475 % and 130 % increase in average YS (1455MPa) and UTS (1557MPa), respectively. Meanwhile, an increase of microhardness was resulted by a factor of three. 2. Evolution of the nanostructure of the CRed Ni-based alloy during annealing and corresponding change in mechanical properties were systematically investigated. Peak YS and UTS of 2047.5 and 2142 MPa, respectively, combined with the microhardness over 600 Hv were obtained when annealing for 5 h at 600 oC. The high strength is due to the formation of the nanoscaled γ′ precipitates and ultrafine grained matrix. 3. Due to the pinning effect of� the� γ′ precipitates on grain boundaries, grain growth can be significant blocked and thus the temperature threshold for the retention of nanocrystalline structure of this NG CRed-alloy is up to 700 oC (0.62 Tm). The annealedCRed alloys maintained NG (90 nm) structure after 60 min annealing at 700 oC and an ultrafine-grained structure (200 nm) were obtained after 60 min annealing at 800 oC.
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Table 1 Composition of the Ni-base alloy.
Alloy
Mn
Si
W
Mo
Al
Ti
C
B
Zr
Cr
Fe
Ni
0.6
0.1
1.0
0.2
1.3~1.9
1.2~1.6
0.03
0.004
0.02
20~25
20
Bal.
composition Content (wt. %)
Table 2 Tensile properties (average value) of the Ni-base alloy with different processing conditions. YS (MPa)
States
UTS (MPa)
A (%)
Solution-treated (STed)
253
684
44
STed, 700℃/1h
454.5
898
41.1
Cold rolled (CRed)
1455
1557
1.95
CRed, 500℃/1h
1907
1949
1.14
CRed, 500℃/5h
1986
2040
1.45
CRed, 600℃/1h
2013
2061
1.4
CRed, 600℃/5h
2047.5
2142
1.94
CRed, 700℃/10min
1939
2007.5
1.81
CRed, 700℃/30min
1844
1923
1.95
CRed, 700℃/1h
1827.5
1951
1.98
CRed, 700℃/5h
1837
1927
2.93
19
Fig. 1. Optical micrograph of Ni-based alloy in solution treatment state before cold
★
★
(b)As-CRed
:γ
（ 222）
★
（ 311）
★
（ 200）
（ 111）
Intensity, a.u.
★
（ 220）
rolling process.
★
(a)Before rolling
20
30
40
50
60
70
80
90
100
2 degree
Fig. 2. XRD patterns of CRed sample and solution treated sample prepared for rolling.
20
（ 220）
: γ' ★
℃
(d)700
65
★
66
81
82
:γ
83
★
（ 222）
64
（ 311）
★
63
（ 221）
37
（ 211）
(e)800
36
（ 200）
35
（ 111）
34
（ 110）
Intensity, a.u.
★
★
℃
(c)600
℃
(b)500
℃
(a)As-CRed 20
40
60
2 degree
80
100
Fig. 3. XRD patterns of CRed-sample and CRed samples subsequently annealed at different temperatures (500, 600, 700, 800 oC) for 1 h. (a)
(b)
21
(c)
(d)
Fig. 4. TEM micrograph of CRed sample processed from microcrystalline Ni-based alloy in solution-treated state. (a) bright field (BF) image; (b) SAED pattern showing the existence of single γ phase; (c) (d) dark field (DF) image corresponding to (111) and (200) crystal plane of BF image in (a), respectively. Average grain size measured from BF is about 50 nm. (a)
(b)
22
(c)
Fig. 5. TEM micrograph of the CRed-sample after annealing at 700 oC for 1 h. (a) bright field (BF) image; (b) dark field (DF) image corresponding to (111) crystal plane of BF image in (a). (c) SAED pattern showing the coexistence of nano-grained γ and γ′ phases; Average grain size measured from BF is about 90 nm. (a)
(b)
23
(c)
Fig. 6. TEM micrograph of the CRed-sample after annealing at 800 oC for 1 h. (a) bright field (BF) image; (b) dark field (DF) image corresponding to (111) crystal plane of BF image in (a). (c) SAED pattern showing the coexistence of nano-grained γ and γ′ phases; Average grain size measured from BF is about 200 nm. 650 600
HV0.5(Kgf.mm-2)
550 500 450 400 350 300 0
200
400
600
800
1000
Annealing temperature,
℃
Fig. 7. Variation of the microhardness of CRed Ni-base alloy with increasing annealing temperature. The annealing time was fixed to be 1 h and the annealing temperature was 500, 600, 700, 800, 900 oC, respectively. The microhardness of the as-deformed samples 24
was also included. 520
HV0.5(Kgf.mm-2)
480
440
400
360
320 0
20
40
60
80
100
Annealing time, h Fig. 8. Variation of the microhardness of CRed Ni-base alloy with increasing annealing time. The annealing temperature was fixed to be 800 oC and the annealing time was 1, 2, 5, 10, 20, 50, 100 h, respectively. The microhardness of the as-deformed samples was also included. 500 600 700
℃
℃
640
-2
HV0.5(Kgf.mm )
℃
600
560
520
480
440 0
100
200
300
400
500
600
Annealing time, min
Fig. 9. Variation of the microhardness of CRed Ni-base alloy with increasing annealing 25
time. The annealing temperature was fixed to 500, 600, 700 oC, respectively. The annealing time in all curves was 5, 10, 20, 30, 60, 120, 180, 300, 600 min, respectively. The microhardness of the as-deformed samples was also included. 650
HV0.5(Kgf.mm-2)
600
550
500
450
400 0
50
100
150
200
250
300
Annealing time, h Fig. 10. Variation of the microhardness of CRed Ni-base alloy with increasing annealing time. The annealing temperature was fixed to 700 oC and the annealing time was 1, 5, 10, 20, 50, 100, 150, 200, 250, 300 h, respectively. The microhardness of the as-deformed samples was also included.
26
2500
(a)As-CRed (b)CR+500 /1h (c)CR+600 /1h (d)CR+700 /1h (e)CR+700 /5h (f)ST+700 /1h (g)ST
Engineering stress, MPa
℃
℃
2000
℃
℃
(b) (c)
(d)
(e)
℃
1500
(a) 1000
(f) 500
(g) 0 0
1
2
3
40
42
44
46
48
Engineering strain, %
Fig. 11. Stress-Strain curves of the samples in different processing conditions.
27
PII: DOI: Reference:
S0921-5093(15)30098-8 http://dx.doi.org/10.1016/j.msea.2015.06.043 MSA32481
To appear in: Materials Science & Engineering A Received date: 26 April 2015 Revised date: 12 June 2015 Accepted date: 13 June 2015 Cite this article as: Yanle Sun, Songqian Xu and Aidang Shan, Effects of annealing on microstructure and mechanical properties of nano-grained Ni-based alloy produced by severe cold rolling, Materials Science & Engineering A, http://dx.doi.org/10.1016/j.msea.2015.06.043 This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting galley proof before it is published in its final citable form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
Effects of Annealing on Microstructure and Mechanical Properties of Nano-grained Ni-based Alloy Produced by Severe Cold Rolling
Yanle Sun1, Songqian Xu2, Aidang Shan1*
1. School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai, 200240, People’s Republic of China 2. Baosteel Special Steel Co., Ltd. Shanghai, 200940, People’ Republic of China
Corresponding author
Tel.: +86 21 54747556; fax: +86 21 54740825 E-mail: [email protected] (Aidang Shan)
1
Abstract Nano-grained (NG) Ni-based alloy with average grain size about 50 nm was produced via severe cold rolling at room temperature. Compared to the solution-treated alloy prepared for rolling, the cold-rolled (CRed) NG alloy exhibited approximately 475 %, 130 % and 180 % increase in the average yield strength (YS), ultimate tensile strength (UTS) and microhardness, respectively. Effect of the annealing on mechanical properties and microstructure of this NG Ni-based alloy were investigated systematically. The annealed-CRed alloys possessed NG (90 nm) and ultrafine-grained structure (200 nm) after 60 min annealing at 700 and 800 oC, respectively, suggesting high thermal stability of the nanostructure. Peak YS and UTS of 2047.5 and 2142 MPa, respectively, were obtained with appropriate annealing temperature and time. The mechanisms of variations in mechanical properties were analyzed based on phase composition and microstructure evolution of the annealed NG Ni-based alloy. Both the high thermal stability and strength are due to the formation of the γ′ precipitates and slight grain growth of the NG matrix.
Key words: Nano-grained (NG); Ni-based alloy; Severe cold-rolling; Mechanical properties; Microstructure; Thermal stability
2
1. Introduction Bulk nano-grained (NG) metals, having a grain size in the nanoscale (10-100 nm), attract considerable attention due to their exceptional physical, chemical and mechanical properties relative to their coarse-grained counterparts [1-4]. Severe plastic deformation (SPD) offers an attractive route to fabricate bulk NG materials. The microstructure and mechanical properties of NG and UFG Ni-based alloy prepared by different SPD processes, such as high-pressure torsion (HPT) [5-7], multiple forging (MF) [5-7] and other special SPD processing [8-10], have been investigated by some researchers. In comparison with the ECAP and HPT, severe cold-rolling process could overcome the major limitations of low productivity of the former and small work-piece size of the latter [11, 12]. While there were some research on the NG pure nickel [13, 14] and single phase Ni-based alloy [15] obtained by severe cold-rolling, there were no study on the microstructure and mechanical properties of a dual-phase NG Ni-based alloy prepared by cold-rolling. Hence, there is a strong scientific incentive to systematically study the evolution of microstructure, phase constitutions and variation of mechanical properties of cold-rolled (CRed) dual-phase NG Ni-based alloy during annealing. The Ni-based alloy in present research consists of a solid solution γ matrix strengthened by stable nanoscaled spherical γ′ phase (Ni3 (Al,Ti)), ordered face centered cubic L12 structure [16]. Our research is motivated by two objectives: first, to use severe cold rolling process to fabricate an NG Ni-based alloy; and second, to systematically investigate the effect of annealing on mechanical properties, phase composition and microstructure of this NG Nibased alloy. The grain size, crystal preferred orientation, phase composition, microhardness as well as tensile properties of the NG alloy were characterized after 3
different annealing process and a high-strength, thermally stable dual-phase nanostructured Ni-based alloy through precipitation were obtained. 2. Experimental procedures The starting material was bulk microcrystalline (20~30 µm, seen in Fig. 1) Ni-based alloy with dimensions of 18 × 15× 100 mm3 and the alloy compositions were given in Table 1. Thin sheet samples of NG alloy with a thickness reduction of 98 % from 18 mm to 0.3 mm was produced by severe cold rolling at room temperature. Prior to this, the bulk materials prepared for cold rolling were solution treated (STed) at 1100 oC for 2 h to ensure that all the initial precipitations were dissolved and a homogeneous microstructure was created. Phase constitutions of the STed, CRed and annealed alloys were determined with Xray diffraction (XRD) technique, using Cu K (𝜆�� = 0.1540562 nm) radiation on a Shimadzu XRD-6000 diffraction instrument, operating at 40 kV and 30 mA between 20 and 100 deg (2θ) at a step of 0.02 ° and a counting time of 0.6 s per step. Transmission electron microscope (TEM) and selected area electron diffraction (SAED) studies were performed using a JEM-2100F operating at 200 kV. The observed sections of CRed and annealed samples for TEM were ND-TD plane. Foils were cut, punched to 3-mm-diameter discs, mechanically polished to a thickness of 50 μm and finally twin-jet polished using an electrolyte of 5 % perchloric acid and 95 % alcohol that was cooled to 243 K with an applied voltage of 40 V. RD, TD and ND stand for rolling, transverse and normal directions, respectively. Specimens prepared for microhardness tests were machined from the CRed-sheets and subsequently annealed at temperatures from 500 to 900 oC for different length of 4
time. Microhardness test was carried out on the ND-RD plane with a load of 500 g and a dwell time of 15 s using the Buehler Micromet Hardness Tester with a Vickers diamond pyramid indenter. For each measured plane, at least 9 points were measured to calculate the mean value of microhardness and the measurement error was no more than 2.5 %. Flat specimens for tensile testing, with a thickness of 0.3 mm, were sectioned along the longitudinal rolling direction from the CRed-sheets by electrical discharge machining. Thickness of tensile specimens cut from the STed bulk material which was prepared for cold rolling was 1 mm. The gage length and width of all the tensile specimens were 15 and 3 mm, respectively. All the tensile tests were conducted at ambient temperature by using a Shimadzu AG-10kNA tension tester at a strain rate of 10-3 s-1. 3. Results 3.1 Phase components and microstructure 3.1.1 XRD analyses Fig. 2 shows the XRD patterns of the STed Ni-based alloy prepared for cold rolling as well as the as-CRed state. All the diffraction peaks corresponded well with those of a pure nickel standard (JCPDS card 01-087-0712), which confirms that the STed Ni-based alloy solely consisted of γ phase and no phase transformation occurred during subsequent cold-rolling deformation. The intensity ratios of all the diffraction peaks in the STed alloy were well in accordance with those of the standard. However, the intensity ratios of (2 2 0) peak of the CRed alloy was much larger than those of STed sample, suggesting preferred orientations along the (2 2 0) crystal plane in the CRed material. Additionally, a slight peak shifting from the pure nickel position to smaller Bragg angles was present in the XRD patterns [17], which was consistent with the fact that both the Ni-based alloys 5
were in solid-solution state. Interestingly, the preferred orientation along the (2 2 0) crystal plane slightly increased with the increase of annealing temperature as illustrated in Fig. 3. From Fig. 3 (b)-(e), new diffraction peaks at 2 of 35.5°, 63.7° and 80.4° were present, which were corresponding to the (1 1 0), (2 1 1) and (2 2 1) crystal planes of γ′ phase. It can be ascertained that the second phase of γ′ precipitated after annealing for 1 h at 500 , 600, 700, 800 oC, respectively. The intensity of (1 1 0), (2 1 1) and (2 2 1) peaks of γ′ phase became stronger with the annealing temperature increasing from 500 to 700 oC and then reduced with the temperature further increasing to 800 oC. The variation of intensity of diffraction peaks of γ′ phase showed that the volume fracture of γ′ precipitation increased gradually as the annealing temperature increasing from 500 to 700 oC and then significantly decreased as the annealing temperature further increasing to 800 oC. 3.1.2 TEM analyses TEM analyses prove that cold rolling process of the Ni-based alloy resulted in complete refinement of the initial coarse-grained structures into nanoscale ones. It is evidenced in Fig. 4 by the bright- field (BF) and dark-field (DF) images as well as the diffraction rings. The TEM images clearly show that the most common feature in the asCRed sample is a random distribution of grains with average size approximately of 50 nm, as measured in Fig. 4 c and d. The microstructure is composed primarily of equiaxed grains, along with some distorted structures. Inspection of BF and DF images shows the presence of high densities of dislocations and that the grain boundaries tend to be curved, irregular and/or ill-defined, which reveals that the grain boundaries introduced by the severe cold-rolling were severely distorted and in a high-energy non-equilibrium 6
configuration. In addition, the SAED pattern in Fig. 4 b exhibits typically continuous rings ascribed to the NG γ phase, which is electron diffraction characteristic of ultrafine grains separated by boundaries having high-angle of misorientation due to the transformation from dislocation cell and sub-grain boundaries as increasing strain induced into the sample [18]. All of these characteristics are typical of materials produced by SPD [19-21]. The CRed samples after annealing for 1 h at 700 and 800 oC, were respectively examined in the TEM to track the evolution of the nano-grains produced by severe cold rolling. Fig. 5 presents the TEM images of CRed samples after annealing at 700 oC for 1 h. The average grain size of the γ matrix is nearly 90 nm, as measured in Fig. 5 b, suggesting a slight increase of the matrix grains in compared with that of the as-CRed material. It is evident that most grains are equiaxed shapes, but with sharp grain boundaries and the morphology of the grains is similar to that of the as-CRed sample. However, grain boundaries in the samples after such an annealing become more distinct, which ascribes to the decrease in the number of defects in near-boundary regions and some extent of stress relaxation. SAED pattern clearly shows the coexistence of� γ and γ′ phases, which is consistent with the XRD analysis in “XRD analyses” section. Additionally, vast scale dislocations can be observed in the matrix and the SAED pattern remains to be continuous rings ascribed to the nano-grains. Fig. 6 presents the TEM images of CRed samples after annealing at 800 oC for 1 h. The average grain size of γ matrix significantly increases beyond the regime of nanoscale and is around 200 nm, as measured in Fig. 6 b. Due to the grain growth, the SAED pattern starts to present separate spots, which further confirms that the grains have grown 7
up apparently. Nevertheless, the grain size distribution emerges a bimodal distribution, as there are several grains below 100 nm. Fig. 6 a and b show that the morphology of γ grains are in equiaxed shape with distinct grain boundaries, which is a structure characteristic with lower density of defects on the grain boundary in comparison with the as-CRed sample. In the structure, the present of thin straight grain boundaries with a junction at an angle of 120° exhibits the occurrence of recrystallization. The fringe contrast of several grain boundaries in both BF and DF images also reveals the intense occurrence of recovery in these grains and that their state verges to the equilibrium. The SAED image in Fig. 6 c clearly shows the existence of γ′ phase. 3.2 Mechanical behavior 3.2.1 Microhardness Fig. 7 presents the variation of microhardness of annealed-CRed samples as a function of the annealing temperature while the annealing time was fixed to be 1 h. It can be seen the microhardness of the as-CRed sample was about 480 Hv (kgf·mm-2), which was almost three times of the solution-treated microcrystalline Ni-based alloy before rolling, of which microhardness was about 169 Hv. As shown in Fig. 7, microhardness of samples increases firstly and then decreases as annealing temperature rising from 500 to 900 oC and a peak microhardness up to 610 Hv was attained at 700 oC. When the annealing temperature increased to 800 oC, the microhardness was still slightly higher than that of the as-CRed alloy. However, an abrupt drop of the microhardness from 480 to 330 Hv suggested a rapid softening of the CRed material when annealing at 900 oC. Fig. 8 shows that the gradually decreased of microhardness as the annealing time prolonged from 1 h to 100 h at 800 oC. 8
The tendency of the microhardness that firstly increased sharply and then fell slightly both at 600 and 700 oC with the annealing time. As shown in Fig.9, rapid increase in microhardness was observed during the subsequent annealing. Even annealed for 5 minutes, a dramatic increase in microhardness was resulted and the incremental value of microhardness was much higher at 700 than 600 oC, while approximately increment higher at 600 than 500 oC. The peak microhardness over 610 Hv was achieved when annealed for 1 h at 700 oC, while for 3 h at 600 oC and 10 h at 500 oC. The annealed-CRed Ni-based alloy exhibits an outstanding microhardness in the γ´strengthened Ni-based alloys. The value of microhardness remained to be essential constant above 600 Hv onset 10 minutes till 2 h and hold 570 Hv as annealing time prolonged to 10 h at 700 oC. As the annealing treatment extended to 300 h at 700 oC in Fig. 10, the microhardness was observed to be gradual decrease with the annealing time. Microhardness of 440 Hv was resulted and this value of microhardness was slightly lower than that of the CRed alloy but 38 % greater than that of the peak-aged (320 Hv) alloy in the conventional microcrystalline state. Furthermore, the top value of microhardenss over 600 Hv hold for a longer duration at 600 oC, starting from 2 to 10 h. The high and stable microhardness of the annealed alloys demonstrates high thermal stability under 700 oC. 3.2.2 Tensile behavior The representative engineering stress-strain curves established from samples machined from STed and CRed Ni-based alloys before and after various annealing processing are presented in Fig. 11. The average values of more detailed mechanical properties are summarized in Table 2. In the STed state, the average yield strength (YS), 9
ultimate tensile strength (UTS) and elongation of coarse-grained samples are 253 MPa, 684 MPa and 44 %, respectively. When the STed alloys annealed at 700 ℃ for 1 h, the average YS and UTS are 79.6 % and 31 % higher than those corresponding to the STedstate, respectively, and the elongation is slightly lower (454.5 MPa, 898 MPa and 41.1 %, respectively). In the case of the CRed state, it is evident that the strength was greatly enhanced after cold-rolling processing. The YS and UTS of the NG Ni-based alloy in the CRed state were 1445 MPa and 1557 MPa, respectively, which were respective over five times and two times of the coarse-grained Ni-based alloy in the STed state prepared for cold-rolling process. However, as predicted, the ductility of the alloy decreased largely after cold-rolling [3, 22]. Strengthening through severe cold-rolling process results in a material with a tensile curve that peaks immediately after yielding (curve of as-CRed). Such a trend of strengthening accompanied by a loss of ductility is generally true for various materials processed by SPD ways [14, 21, 22]. Based on the average value of mechanical properties for all samples machined from a processd alloy, the YS and UTS further increased in the annealed-CRed Ni-based alloys. When the CRed alloys were annealed at 600 oC for 1 h, as shown in Fig. 11 and Table. 2, the YS and UTS were 2013 and 2061 MPa, respectively. What is more, a higher YS and UTS up to 2047.5 and 2142 MPa were produced as the annealing time prolonged to 5 h. However, the high strength was accompanied by ductility minimum. The high YS and UTS, together with the microhardness of the annealed-CRed alloys suggested stable mechanical behavior after annealing at 600 oC. When the annealing temperature was fixed at 700 oC, rapid enhancement of the YS and UTS by 33 % and 28.9 % were resulted, respectively, as annealing for 10 min (1939 and 2007.5 MPa, respectively). The 10
YS and UTS were 1837 and 1927 MPa, respectively, as annealing for 5 h. When the annealing time extending from 10 min to 5 h, the YS and UTS are holding above 1800 and 1900 MPa, respectively. However, the elongation is limited and slightly increased with the increase of annealing time. When the factors of the annealing temperature and time was fixed, alternatively, and the other one was considered as variable factor, the nature in change of YS and UTS was quite in accordance with that of the microhardness. 4. Discussion 4.1Microstructure and Mechanical properties of the CRed Ni-based alloy. 4.1 Strengthening mechanism of CRed Ni-based alloy. In the present work, the bulk NG Ni-based alloy with average grain size of 50 nm was fabricated by severe cold-rolling processing. X-ray diffraction (XRD) analysis of asCRed sample suggested that all the additional alloy elements were completely dissolved into the Ni matrix, as only diffraction peaks of the γ phase were detected. Thus the possible strengthening effect of solute atom clusters or precipitates can be neglected. Increase of the microhardness and strength in the as-CRed alloy mainly results from the grain boundary strengthening. High fraction grain boundaries surrounding the nanograins act as obstacles to the movement of the dislocations which would pile up near the grain boundaries, leading to a steep enhancement in the microhardness, YS and UTS [11]. TEM observation revealed that the NG matrix which typically composed of grains around 50 nm were produced by the severe cold-rolling processing. Moreover, the high defect densities inherent in the deformed materials make a further contribution to the strength. Hugnes and Hansen had testified that cold- rolling process of nickel to a strain of 90 % 11
resulted in a greater increase of strength than that calculated in the Hall-Petch rule [13, 23]. Similar results were verified in other materials such as Al alloys by high pressure torsion (HPT) [24]. Since the initial establishment of the Hall-Petch equation was based on the un-deformed materials, it was reasonable that the low angle grain boundaries (LABs) and non-equilibrium state of grain boundaries (NGBs) were not counted, which were only abounded in the severe plastic deformed materials. But in the materials processed by SPD techniques, the NGBs and LABs introduced into the material during processing would play an considerable contribution to the strength [25]. In case of the CRed Ni-based alloy in our research, the contribution of these two boundaries are taken into account and the total strength of the CRed samples can be estimated as followed [26, 27] :
σ y ~ 0 kd -1/2 LAB NGB
(1)
Where 0 is the threshold stress, k is the Hall-Petch constant, d is grain size, kd -1/2 � represents the strengthening contribution of the high angle grain boundaries (HAB) between grains. The LAB and NGB represent the strengthening contribution of the LABs and NGBs, respectively. TEM observation in Fig. 4 reveals that the microstructure is basically composed of HABs, which make a major contribution to the strengthening in the as-CRed samples. From the Eq. (1), it is evident that the ultra-high YS and UTS of the as-CRed Ni-alloy are mainly resulted from the extremely fine grains and additional enhanced by high defect densities [28]. In the case of the microhardness, the Hall-Petch relationship in Eq. (1) can be applied and its values significantly increased from 169 Hv of the coarse-grained STed alloy to 480 Hv of the NG CRed-alloy. 12
4.2 Improved strength of the NG Ni-based alloy via annealing When the as-CRed Ni-base samples annealed at 500, 600 and 700 oC, further increase of the YS, UTS and microhardness were resulted. No doubt the very rapid hardening is mainly due to the rapid formation of the nanoscaled γ′ precipitates in the NG matrix while grain growth of the NG matrix is slight at these temperatures. This is confirmed by the TEM observations in Fig. 5, “XRD analysis” in Fig. 3 and the microhardness variation in Fig. 9. Higher microhardness and strength correlate with higher volume fraction of the γ′ precipitates. While we have not explicitly imaged the evolution of the nanoscaled precipitates in the alloy, such rapid hardening in nanostructured precipitation-strengthening alloy is known to be due to the rapid formation of the second-phase precipitates [9, 29]. Microhardness of the CRed alloy increased from 480 to 615 Hv while the corresponding YS and UTS increased from 1455 and 1557 to 2047.5 and 2142 MPa, respectively. Generally, the correlation between the YS and microhardness has been determined to be [30, 31]:
Hv c y
(2)
Where y is the uniaxial YS and c is the elastic constraint factor. For blunt pyramid indenters, it was reported that c 0.3 when the hardness and yield strength are measured in kgf·mm-2 and MPa [30, 31] . According to calculation of microhardness values on the ND-RD plane of the as-CRed and annealed-CRed samples, the factor of c was slightly waved for the different states of material. Since both the microhardness and mechanical strength are significantly affected by microstructure, the correlation between the YS and microhardness are also affected by microstructure [32]. In the as-CRed alloy, the high 13
density of defects such as dislocations were expected to affect this correlation and led to a large value of factor to 0.33. When the as-CRed alloy annealed in various conditions, a slight fluctuation of the value of factor c was resulted for the evolution of microstructure during annealing. According to the SAED, TEM and XRD analysis, such annealed-CRed alloys were composed of the NG matrix of γ phase, typically embedded with dispersed γ′ precipitation. In this precipitate-strengthened NG Ni-based alloy, strengthening is due to the stable ultra-fine γ grains and the dispersed of nanoscaled γ′ precipitates. Thus, the effective strengthening mechanism are the Hall-Petch strengthening due to the NG matrix and the Orowan mechanism due to the spherical γ′ precipitates. Thus, resulting strengthen (τ y ) in the annealed-CRed Ni-based alloy can be approximately expressed as followed [9]:
τy y
Gb
λ
(3)
Where y �is the strength of the as-CRed alloy showed in Eq. (1),� G is the shear modulus, b �is the Burgers vector and λ �is the mean spacing between γ′ particles. The part of��
Gb
λ
in Eq. (3) represents the effect of Orowan strengthening attributed to the γ′
particles, which suggest that dislocations circumvent hard particles leaving dislocation loops behind. Since d 1/2 �(in� Eq. (2))� and� λ �(in� Eq. (3))� are in an order of magnitude, these two strengthening effects are comparable and either one can be neglected. Thus, the resulting strength is the tradeoff between the effect of precipitation strengthening and matrix softening. When the CRed alloy annealed at 800 oC for 1 h, recrystallization of the 14
structure occurred as the present of distinctly straight grain boundaries with a junction at an angle of 120° and the grains significantly grown to about 200 nm. The softening effect of the recrystallization and grain growth prevailed over the strengthening effect of the precipitation as the annealing time prolonged over 2 h at this temperature. Thus the strength was observed to be decrease. From the Orowan formula (
Gb
λ
), it can be seen that the finer the precipitate size
and the larger their volume fraction, the greater effect of precipitation strengthening produced. In the “XRD analysis” section, it is clear that the intensities of diffraction peaks of γ′ phase become much stronger as annealing temperature increasing from 500 to 700 oC, showing evident increase in the volume fraction of γ′ precipitation. Foreparts of curves (forward the value of 60 on the horizontal axis) in Fig. 9 suggested that increase of annealing temperature from 500 to 700 oC would accelerate the rapid hardening of the γ′ precipitates. Evolution of microhardness of annealed-CRed samples was in accordance with the XRD results. When the CRed alloy annealed under 800 oC, strength softening of the NG matrix were more apparent for the growth of grains and annihilation of a part of LABs and NGBs and the strengthening effect of nanoscaled γ′ precipitates would go down for less volume fraction of the precipitation. It is reported that a peak hardness value of 630 Hv was achieved in the γ″ and γ′strengthened NG� Inconel 718 and the microhardness of the counterpart with coarse grains after solution treatment was 300 Hv [9]. In the present work, a peak value of microhardness of this γ′ precipitate-strengthened NG Ni-based alloy was 615 Hv, almost by a factor of four higher than that of the coarse-grained alloy in the STed state, which 15
was about 169 Hv. Comparison of the above results, it can be seen much more significant enhancement of the microhardness was achieved in this Ni-based alloy. 4.3 High thermal stability In the present work, a notable thermal stability of the NG CRed-alloy was observed. Most recently study shows that the critical temperature of nanostructured Inconel 718 with grain size of 100 nm produced by plane strain machining is 0.6 Tm (Tm, melting point in K) [9]. It has been reported that NG Ni-20%Cr and Inconel 718 produced by HPT, typically with a mean grain size of 50-80 nm, are thermally stable up to 0.46 and 0.5 Tm, respectively [7]. The temperature threshold for the retention of nanocrystalline structure of the Inconel 718 is higher than that of the Ni-20%Cr, which is due to the effect of the uniformly distributed coherent disk-type-γ′′-precipitates within grains in the Inconel 718. Improvement of thermal stability of the nanocrystalline alloy by introducing a dense dispersion of fine second-phase precipitates has also been reported in many other alloys such as Ti alloys and Al alloys [29, 33-35]. In the present research, the Tm� of the Ni-based alloy is 1300 oC (1573K) that calculated by the JMatPro software. TEM observation of the CRed and annealed-CRed alloys shows that the average grain size maintains below 100 nm when annealed for 60 min at 700 oC and thus this NG matrix with grain size about 50 nm is stable up to 0.62 Tm. In the view of above discussion, the high thermal stability is closely associated with the pinning effect of the� γ′ precipitates on grain boundaries and thus grain growth can be significantly blocked. Further research on the thermal stability of NG Ni-based alloy is undertaken. 5. Conclusion 1. NG Ni-based alloy with average grain size of about 50 nm was prepared via 16
severe cold-rolling at room temperature. As compared to the STed alloy before rolling, the CRed alloy exhibits approximately 475 % and 130 % increase in average YS (1455MPa) and UTS (1557MPa), respectively. Meanwhile, an increase of microhardness was resulted by a factor of three. 2. Evolution of the nanostructure of the CRed Ni-based alloy during annealing and corresponding change in mechanical properties were systematically investigated. Peak YS and UTS of 2047.5 and 2142 MPa, respectively, combined with the microhardness over 600 Hv were obtained when annealing for 5 h at 600 oC. The high strength is due to the formation of the nanoscaled γ′ precipitates and ultrafine grained matrix. 3. Due to the pinning effect of� the� γ′ precipitates on grain boundaries, grain growth can be significant blocked and thus the temperature threshold for the retention of nanocrystalline structure of this NG CRed-alloy is up to 700 oC (0.62 Tm). The annealedCRed alloys maintained NG (90 nm) structure after 60 min annealing at 700 oC and an ultrafine-grained structure (200 nm) were obtained after 60 min annealing at 800 oC.
17
References [1] R.Z. Valiev, R.K. Islamgaliev, I.V. Alexandrov, Progress in materials science, 45 (2000) 103-189. [2] H. Gleiter, progress in Materials Science, 33 (1989) 223-315. [3] Y.M. Wang, M.W. Chen, F.H. Zhou, E. Ma, Nature, 419 (2002) 912-915. [4] L. Lu, Y. Shen, X. Chen, L. Qian, K. Lu, Science, 304 (2004) 422-426. [5] X. Sauvage, S. Mukhtarov, IOP Conference Series: Materials Science and Engineering, IOP Publishing, 2014, pp. 012173. [6] S.K. Mukhtarov, V. Valitov, N. Dudova, (2010). [7] S. Mukhtarov, N. Dudova, V. Valitov, Materials Science and Engineering: A, 503 (2009) 181-184. [8] S. Swaminathan, M. Ravi Shankar, B. Rao, W.D. Compton, S. Chandrasekar, A. King, K. Trumble, Journal of Materials Science, 42 (2007) 1529-1541. [9] M. Ravi Shankar, B.C. Rao, S. Chandrasekar, W. Dale Compton, A.H. King, Scripta Materialia, 58 (2008) 675-678. [10] C. Saldana, P. Yang, J.B. Mann, W. Moscoso, D.D. Gill, S. Chandrasekar, K.P. Trumble, Materials Science and Engineering: A, 503 (2009) 172-175. [11] I. Sabirov, M.Y. Murashkin, R.Z. Valiev, Materials Science and Engineering: A, 560 (2013) 1-24. [12] Y. Estrin, A. Vinogradov, Acta Materialia, 61 (2013) 782-817. [13] D.A. Hughes, N. Hansen, Acta Materialia, 48 (2000) 2985-3004. [14] G. Wilde, G.P. Dinda, H. Rösner, Advanced Engineering Materials, 7 (2005) 11-15. [15] P.P. Bhattacharjee, R.K. Ray, N. Tsuji, Acta Materialia, 57 (2009) 2166-2179. [16] Y.Q. Chen, E. Francis, J. Robson, M. Preuss, S.J. Haigh, Acta Materialia, 85 (2015) 199-206. [17] J.C. Villegas, L.L. Shaw, Acta Materialia, 57 (2009) 5782-5795. [18] P. Mckenzie, R. Lapovok, Acta Materialia, 58 (2010) 3198-3211. [19] A. Zhilyaev, A. Gimazov, T. Langdon, Journal of Materials Science, 48 (2013) 4461-4466. [20] Z. Luo, H. Zhang, N. Hansen, K. Lu, Acta Materialia, 60 (2012) 1322-1333. [21] K. Edalati, Z. Horita, Materials Science and Engineering: A, 528 (2011) 7514-7523. [22] R. Valiev, Nat Mater, 3 (2004) 511-516. [23] N. Hansen, Scripta Materialia, 51 (2004) 801-806. [24] R.Z. Valiev, N.A. Enikeev, M.Y. Murashkin, V.U. Kazykhanov, X. Sauvage, Scripta Materialia, 63 (2010) 949-952. [25] R.Z. Valiev, Materials Transactions, 55 (2014) 13-18. [26] N. Krasilnikov, W. Lojkowski, Z. Pakiela, R. Valiev, Materials Science and Engineering: A, 397 (2005) 330-337. [27] R.Z. Valiev, T.G. Langdon, Advanced Engineering Materials, 12 (2010) 677-691. [28] R.Z. Valiev, M.Y. Murashkin, E.V. Bobruk, G.I. Raab, Materials transactions, 50 (2009) 87. [29] M.R. Shankar, S. Chandrasekar, A.H. King, W.D. Compton, Acta Materialia, 53 (2005) 4781-4793. [30] X.C. Liu, H.W. Zhang, K. Lu, Science, 342 (2013) 337-340. [31] C.H. Cáceres, J.R. Griffiths, A.R. Pakdel, C.J. Davidson, Materials Science and Engineering: A, 402 (2005) 258-268. [32] Z. Li, L. Fu, B. Fu, A. Shan, Materials Science and Engineering: A, 558 (2012) 309-318. [33] R. Huang, Y. Han, Journal of Alloys and Compounds, 554 (2013) 1-11. [34] W. Chen, Q. Sun, L. Xiao, J. Sun, Materials Science and Engineering: A, 536 (2012) 223-230. [35] M.R. Shankar, S. Chandrasekar, W.D. Compton, A.H. King, Materials Science and Engineering: A, 410-411 (2005) 364-368.
18
Table 1 Composition of the Ni-base alloy.
Alloy
Mn
Si
W
Mo
Al
Ti
C
B
Zr
Cr
Fe
Ni
0.6
0.1
1.0
0.2
1.3~1.9
1.2~1.6
0.03
0.004
0.02
20~25
20
Bal.
composition Content (wt. %)
Table 2 Tensile properties (average value) of the Ni-base alloy with different processing conditions. YS (MPa)
States
UTS (MPa)
A (%)
Solution-treated (STed)
253
684
44
STed, 700℃/1h
454.5
898
41.1
Cold rolled (CRed)
1455
1557
1.95
CRed, 500℃/1h
1907
1949
1.14
CRed, 500℃/5h
1986
2040
1.45
CRed, 600℃/1h
2013
2061
1.4
CRed, 600℃/5h
2047.5
2142
1.94
CRed, 700℃/10min
1939
2007.5
1.81
CRed, 700℃/30min
1844
1923
1.95
CRed, 700℃/1h
1827.5
1951
1.98
CRed, 700℃/5h
1837
1927
2.93
19
Fig. 1. Optical micrograph of Ni-based alloy in solution treatment state before cold
★
★
(b)As-CRed
:γ
（ 222）
★
（ 311）
★
（ 200）
（ 111）
Intensity, a.u.
★
（ 220）
rolling process.
★
(a)Before rolling
20
30
40
50
60
70
80
90
100
2 degree
Fig. 2. XRD patterns of CRed sample and solution treated sample prepared for rolling.
20
（ 220）
: γ' ★
℃
(d)700
65
★
66
81
82
:γ
83
★
（ 222）
64
（ 311）
★
63
（ 221）
37
（ 211）
(e)800
36
（ 200）
35
（ 111）
34
（ 110）
Intensity, a.u.
★
★
℃
(c)600
℃
(b)500
℃
(a)As-CRed 20
40
60
2 degree
80
100
Fig. 3. XRD patterns of CRed-sample and CRed samples subsequently annealed at different temperatures (500, 600, 700, 800 oC) for 1 h. (a)
(b)
21
(c)
(d)
Fig. 4. TEM micrograph of CRed sample processed from microcrystalline Ni-based alloy in solution-treated state. (a) bright field (BF) image; (b) SAED pattern showing the existence of single γ phase; (c) (d) dark field (DF) image corresponding to (111) and (200) crystal plane of BF image in (a), respectively. Average grain size measured from BF is about 50 nm. (a)
(b)
22
(c)
Fig. 5. TEM micrograph of the CRed-sample after annealing at 700 oC for 1 h. (a) bright field (BF) image; (b) dark field (DF) image corresponding to (111) crystal plane of BF image in (a). (c) SAED pattern showing the coexistence of nano-grained γ and γ′ phases; Average grain size measured from BF is about 90 nm. (a)
(b)
23
(c)
Fig. 6. TEM micrograph of the CRed-sample after annealing at 800 oC for 1 h. (a) bright field (BF) image; (b) dark field (DF) image corresponding to (111) crystal plane of BF image in (a). (c) SAED pattern showing the coexistence of nano-grained γ and γ′ phases; Average grain size measured from BF is about 200 nm. 650 600
HV0.5(Kgf.mm-2)
550 500 450 400 350 300 0
200
400
600
800
1000
Annealing temperature,
℃
Fig. 7. Variation of the microhardness of CRed Ni-base alloy with increasing annealing temperature. The annealing time was fixed to be 1 h and the annealing temperature was 500, 600, 700, 800, 900 oC, respectively. The microhardness of the as-deformed samples 24
was also included. 520
HV0.5(Kgf.mm-2)
480
440
400
360
320 0
20
40
60
80
100
Annealing time, h Fig. 8. Variation of the microhardness of CRed Ni-base alloy with increasing annealing time. The annealing temperature was fixed to be 800 oC and the annealing time was 1, 2, 5, 10, 20, 50, 100 h, respectively. The microhardness of the as-deformed samples was also included. 500 600 700
℃
℃
640
-2
HV0.5(Kgf.mm )
℃
600
560
520
480
440 0
100
200
300
400
500
600
Annealing time, min
Fig. 9. Variation of the microhardness of CRed Ni-base alloy with increasing annealing 25
time. The annealing temperature was fixed to 500, 600, 700 oC, respectively. The annealing time in all curves was 5, 10, 20, 30, 60, 120, 180, 300, 600 min, respectively. The microhardness of the as-deformed samples was also included. 650
HV0.5(Kgf.mm-2)
600
550
500
450
400 0
50
100
150
200
250
300
Annealing time, h Fig. 10. Variation of the microhardness of CRed Ni-base alloy with increasing annealing time. The annealing temperature was fixed to 700 oC and the annealing time was 1, 5, 10, 20, 50, 100, 150, 200, 250, 300 h, respectively. The microhardness of the as-deformed samples was also included.
26
2500
(a)As-CRed (b)CR+500 /1h (c)CR+600 /1h (d)CR+700 /1h (e)CR+700 /5h (f)ST+700 /1h (g)ST
Engineering stress, MPa
℃
℃
2000
℃
℃
(b) (c)
(d)
(e)
℃
1500
(a) 1000
(f) 500
(g) 0 0
1
2
3
40
42
44
46
48
Engineering strain, %
Fig. 11. Stress-Strain curves of the samples in different processing conditions.
27