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Enhanced boronizing kinetics of alloy steel assisted by surface mechanical attrition treatment
Accepted Manuscript Enhanced boronizing kinetics of alloy steel assisted by surface mechanical attrition treatment Haopeng Yang, Xiaochun Wu, Zhe Yang, Shengjun Pu, Hongbin Wang PII: DOI: Reference:
S0925-8388(13)03130-7 http://dx.doi.org/10.1016/j.jallcom.2013.12.151 JALCOM 30215
To appear in: Received Date: Revised Date: Accepted Date:
2 July 2013 16 December 2013 16 December 2013
Please cite this article as: H. Yang, X. Wu, Z. Yang, S. Pu, H. Wang, Enhanced boronizing kinetics of alloy steel assisted by surface mechanical attrition treatment, (2013), doi: http://dx.doi.org/10.1016/j.jallcom.2013.12.151
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Enhanced boronizing kinetics of alloy steel assisted by surface mechanical attrition treatment Haopeng Yang∗, Xiaochun Wu, Zhe Yang, Shengjun Pu, Hongbin Wang School of Material Science and Engineering, Shanghai University, Shanghai 200072, People’s Republic of China
Abstract: A nanostructured surface layer was fabricated on AISI H13 steel by means of surface mechanical attrition treatment(SMAT). Boronizing behaviors of the SMAT samples were systematically investigated in comparison with their coarse-grained counterparts. The boron diffusion depth of the SMAT sample with pack boriding treatment at 600℃ for 2h was about 8µm, which was much deeper than that of the coarse-grained sample. A much thicker borided layer on the SMAT sample can be synthesized by a duplex boronizing treatment at 600℃ followed by at a higher temperature. The borided layer was composed with monophase of Fe2B and the growth of it exhibited a (002) preferred orientation. Moreover, the activation energy of boron diffusion for the SMAT sample is 140.3kJ/mol, which is much lower than 209.4kJ/mol for the coarse-grained counterpart. The results indicate that the boronizing kinetics can be significantly enhanced in the SMAT sample with a duplex boronizing treatment. Furthermore, the thermal fatigue tests show that the borided layer with excellent oxidation resistance and mechanical strength at elevated temperatures could effectively delay the thermal fatigue cracks initiation and impede their propagation. Therefore, the thermal fatigue property of H13 steel with a duplex boronizing treatment can be improved remarkably. ∗ Corresponding author. Tel.: +86 21 56331153; fax: +86 21 56331461. E-mail addresses: [email protected] (H.P.Yang), [email protected] (X.C.Wu). 1
Keywords: H13 steel; Surface mechanical attrition treatment; Diffusion; Boronizing kinetics 1. Introduction AISI H13 steel is one of the most important hot work die steels. In the service process, it is subjected to high loads and high temperatures, which produce serious damages of tool surfaces because of wear, plastic deformation and thermal fatigue. In order to achieve longer service life, various thermochemical surface hardening treatments have been applied to improve its wear performance, oxidation resistance and thermal fatigue property, such as gas and plasma nitriding [1-3], physical vapor deposition of hard ceramic coating [4], packed powder chromizing treatment [5] and duplex treatments [6]. The technology of duplex surface modification has attracted extensive attention in recent years. Torres et al. [4] used nitriding as pretreatment to prepare TiAlN/TiN coating with different thicknesses on H13 substrate, which could improve coating-substrate adhesion, and the wear rate of the duplex coatings can be decreased significantly. However, duplex treatments usually demand the use of rather complicated expensive apparatus for plasma nitriding and physical or chemical vapor deposition. These disadvantages hinder their wide applications. Boronizing is an effective surface hardening process. The borided layer possesses excellent adhesion to the substrate due to the nature of the thermal diffusion process. The boronizing technique can be carried out in solid, liquid or gaseous medium [7]. Among the various boronizing processes, solid-state pack boriding treatment is the most frequently used. Genel et al. [8] prepared both phases of FeB and Fe2B on H13 2
steel in the solid medium consisting of Ekabor-I powders at 800, 900, 1000℃for periods of 1–5h. Taktak et al. [9] reported the FeB and Fe2B can be formed on H13 steels in slurry salt bath consisting of borax, boric acid and ferrosilicon at temperature range of 800–950℃ for 3, 5 and 7h. However, the disadvantages of pack boriding treatment are the requirements of relatively high processing temperature (850–1000℃) and long process duration to obtain a useful borided layer thickness. Studies have been carried out to improve the efficiency of pack boriding treatment over the past decades. It is well known that borided layer generated in thermochemical treatment depends on boriding condition and on the property of the material itself. Both factors are strongly affected by grain boundaries and defect density. The combination of pack boriding treatment with surface mechanical attrition treatment(SMAT) could provide an alternative approach to observably improve the efficiency of boronizing kinetics. After SMAT, various metals, such as pure Fe [10], 38CrMoAl [11], AISI 321 austenitic stainless [12] and other alloys [13] possess a nanostructured surface layer. Due to the significant enhanced diffusion and chemical reaction kinetics in the formed nanostructured surface layer, the hardened diffusion layer with disperse compounds on the substrate has been fabricated on several ferrous alloys after the subsequent gas nitriding [10,11] or packed powder chromizing treatment [14,15]. In addition, the efficiency of thermochemical treatment has been increased evidently relative to that of the conventional one without SMAT. Here, we report that pack boriding treatment was carried out for the hot work die 3
steel H13 assisted by SMAT. The boronizing kinetics in the nanostructured surface layer fabricated by SMAT was studied. Furthermore, the thermal fatigue property of H13 steel with a duplex boronizing treatment was investigated by means of thermal fatigue tests with 1000, 2000 and 3000 continuous cycles from 20 to 700℃. 2. Experimental procedure 2.1 Sample preparation The chemical compositions of AISI H13 steel used in the experiments contain (wt.%) 0.42C, 4.93Cr, 1.40Mo, 0.98Si, 0.87V, 0.30Mn, 0.018P, 0.005S and balance Fe. Before SMAT, the ferritic steel with the shape(80×80×4mm) was mirror polished. The set-up and procedures of SMAT treatment had been described in a previous paper [16]. In our experiments, the SMAT to H13 steels was performed under vacuum at room temperature for 60min with a vibrating frequency of 50Hz using bearing steel balls with 8mm diameter. Because the sample surface was plastically deformed with high strains and high strain rates, grains in the surface layer are effectively refined. The small samples with dimensions of 15×15×4mm were machined from the bulk sample mentioned above and ultrasonically cleaned in acetone. The pack boriding treatment was as follows: the constituents of the boriding media were B4C(10wt.%), KBF4(5wt.%), charcoal(5wt.%) with the balance of SiC. The SMAT samples and the untreated ones were packed in the powder mix and sealed in a stainless steel container. Boronizing was performed in an electrical resistance furnace at 600℃(T1) for 2h followed by heating at different higher temperatures(T2) for various time. In order to stabilize the nanostructures in the SMAT surface layer, 4
boronizing was carried out at T1, which was useful to form a much thicker borided layer on H13 steel with a subsequent boronizing process at a higher temperature T2. The boronizing temperatures were monitored by a NiCr–NiSi thermocouple with an accuracy of ±2℃. All samples were cooled with furnace. Finally, the borided samples were quenched at 1030℃ and tempered twice at 580℃ in vacuum oven. 2.2 Thermal fatigue tests Thermal fatigue tests were carried out in a high frequency induction test system. The schematic diagram of the thermal fatigue test system is shown in Fig.1. The Uddeholm self-restricted method was used for the thermal fatigue tests. The size and geometry of a thermal fatigue test specimen are shown in Fig.2(a). The integral specimen is composed with two cylinders. The diameter of one cylinder is 5mm and the other is 10mm.Two study areas were prepared symmetrically on the bigger cylinder, as indicated by arrows in Fig.2(b). One study area was the polished surface, the other was the borided surface assisted by SMAT. An induction heating system and a water shower system were used for heating and cooling, respectively. The thermal fatigue tests were carried out from 20 to 700℃.The error of the temperature was about 10℃, considering the fast rates of heating and cooling. The specimen was heated from 20 to 700℃ in 3 seconds and was kept for 1 second to allow the specimen to expand and to reach the equilibrium. Then, the specimen was cooled with water shower for 8 seconds and the temperature could be decreased to about 20℃ rapidly. The total time of one cycle was 12 seconds. The specimen was subjected to repeated compressive and tensile stress. The number of cycles in our experiments 5
included 1000, 2000 and 3000. After thermal fatigue test, the specimen was soaked in HCl solution(10% vol.%) and the oxide skin on specimen can be removed. The crack surfaces were observed with optical microscope to identify crack mode. Moreover, the morphologies of cracks formed in the region beneath the crack surfaces were studied after sectioning the cracked specimen longitudinally. 2.3 Characterization The
X-ray
diffraction(XRD)
measurements
were
obtained
with
a
RigakuD/Max-RBX-ray diffractometer by using CuKa(40kV,40mA) radiation and a secondary beam graphite monochromator. Cross-sectional observations of the borided samples were performed on a JSM-6301F scanning electron microscope(SEM). Microstructural features in the surface layer were characterized by using a field emission transmission electron microscopy(TEM,JEOLJEM-2010F) equipped with an energy dispersive spectral(EDS) analyzer. Thin foil samples for TEM observations were prepared by cutting, grinding and dimpling with a final ion thinning at low temperatures. Microstructures of cross section of the borided layers etched in 4% nital solution were examined by an optical microscope. The hardness distributions of the cross-sectioned borided layers were measured by a Vickers microhardness tester with a load of 100g force. A LECO GDS850A glow-discharge optical emission spectrometer(GD-OES) was used to measure the concentrations of boron element in the surface layers of the borided samples. 3. Results and discussion 3.1 Microstructure of the SMAT sample 6
Surface mechanical attrition is an effective technique to realize the surface self-nanocrystallization on metallic materials. Fig.3(a) and (b) show that the TEM observations in the top surface layer treated by SMAT. Due to the very high strain and strain rate(102 to 103s−1) [16], extremely fine equiaxed ferrite grains with random crystallographic orientations are formed, as revealed by the corresponding selected-area electron diffraction pattern. Carbide phase may be progressively refined into smaller particles or dissolved into the ferrite phase [14,17]. Therefore, no carbides diffraction ring can be detected. Fig.3(c) displays the statistical distribution of grain size. The size of vast majority of grains is at 10-15nm. The grain size increases gradually along the depth in the surface layer. The similar experimental results were also obtained by some previous reports [5,18]. When comparing to the conventional grain boundaries(GBs) in the annealed coarse-grained samples, the GBs formed by the severe plastic deformation are usually in non-equilibrium states which are associated with a high density of defects or a high stored energy [19]. GBs diffusivity in the deformation layer produced by SMAT [20] or equal-channel angular pressing [21,22] will be considerably enhanced relative to that along high angle GBs in the coarse-grained counterpart. It should be noted that the thermal stability of the nanostructured surface layer can be preserved up to 600℃ according to the previous study results [5]. 3.2 Diffusion of boron in the nanostructured surface layer The optical micrographs of the cross section of the SMAT sample and the coarse-grained counterpart with pack boriding treatment at 600℃ for 2h are showed 7
in Fig. 4(a) and (b), respectively. The microstructural morphology in the surface layer with dark contrast region differs from that in the internal substrate for SMAT sample and the boundary between them is distinctive, which indicate that the chemical components and the crystal structures of the two regions are different, resulting in the different corrosion potential for the top surface layer and the internal substrate etched in nital solution. However, the similar result does not obviously exist in the coarse-grained sample. The concentration of boron along the depth from the treated surface of the SMAT sample and the coarse-grained counterpart are presented in Fig. 4(c), which is determined by GD-OES tests. The peak value of concentration in the top surface layer for the SMAT sample is about 6 wt.%, while the peak value for the coarse-grained sample is only about 0.6 wt.%. At the same time, the depth of boron diffusion in SMAT sample is about 8µm, which is much deeper than that in the coarse-grained sample. Based on the experimental results mentioned above, it can be concluded that the thermal diffusion of boron in the nanostructured surface layer fabricated by SMAT can be evidently enhanced when compared to that in the coarse-grained sample with the same experimental parameter. According to the diffusion equation for semi-infinite solid diffusion [23]:
X = α Dmt
(1)
where α is a constant (about 1 in the present case), t is the duration of the pack boriding treatment, and Dm is the mean diffusion coefficient. The measured diffusion depth into the SMAT sample is larger than that into the coarse-grained counterpart, 8
indicating the change of diffusion mechanism in the SMAT surface layer. Various non-equilibrium GBs, along which the diffusion rate of boron is much higher than that along the relaxed GBs in the coarse-grained counterpart, dominate the much enhanced boronizing kinetics in the SMAT surface layer at 600℃. The mean GBs diffusion coefficient calculated from the measured boron diffusion depth in the SMAT sample according to Eq. (1) is 8.9×10-15 m2/s, which is much higher than 1.3×10-15 m2/s for the coarse-grained sample. 3.3 Microstructure and hardness of the borided layer The pack boriding treatment at such a low temperature was found to facilitate stabilizing the nanostructures in the SMAT surface layer, which was useful to form a much thicker borided layer to enhance the surface properties of tool steels with a subsequent boronizing process at a higher temperature. Fig. 5(a) and (b) show the cross-sectional SEM observations for the coarse-grained sample and the SMAT sample after a duplex boronizing treatment(DBT) at T1 =600℃ for 2h followed by at T2 =850℃ for 8h, respectively. It can be seen that the continuous borided layers are on the two sample surfaces. The thicknesses of borided layers on the SMAT sample and the coarse-grained counterpart are 40µm and 30µm, respectively. Neither FeB(orthorhombic) nor Fe2B(tetragonal) has cubic crystal symmetry [24]. Although H13 steel possesses a certain amount of alloying elements, boron diffusion in the borided layer is of anisotropic nature at high temperatures. Therefore, the morphologies of coating layers exhibit tooth-like growth into substrate. The microhardness gradients for the two samples are showed in Fig. 5(c). For both samples, the graphs include three distinct regions, which are borided layer, transition zone and matrix that is not affected by boron diffusion. The microhardness of borides formed on the surface of SMAT sample is about 1800HV0.1 and it is roughly 3.5 times harder than that of the substrate. In comparison, the microhardness of borides is about 1600HV0.1 for the coarse-grained counterpart. Owing to significant refined grain sizes 9
in the nanostructured surface layer of SMAT sample, the refined grains of borides on H13 steel could be synthesized with a DBT. However, after the borided samples were quenched at 1030℃ and tempered twice at 580℃, the refined grains of borides on SMAT sample could grow into larger crystalline grains, and the difference between the grain sizes of borides on SMAT sample and that on coarse-grained counterpart is not significantly obvious. At the same time, the continuous and dense microscopic structures of borides on them are not evidently different. Therefore, there is not a substantial difference in the microhardness gradients between the SMAT and coarse-grained samples with a DBT. In addition, the microhardness of transition zone is a little softer than that of the substrate. Similar results were reported for H11, H12 and H13 steels [25,26]. The decrease in hardness can be ascribed to concentration of the silicon atom in front of borided layer, as silicon is not soluble in it. It is well known that silicon is a ferrite-forming element and the high concentration of silicon ahead of the borided layer suggests that this region remained in ferrite at the austenitizing temperature and consequently did not subsequently transform into martensite as did in the internal substrate during the quenching process [25]. As shown in Fig. 6, based on the pack boriding treatment at T1=600℃ for 2h, a much thicker borided layer was formed on the SMAT sample than that on the coarse-grained counterpart after the boronizing treatment at any T2 from 750-850℃ for 8h. The thicknesses in both samples increase with an increasing T2. It can be seen that the low temperature treatment at 600℃ will enhance the boronizing kinetics of the SMAT sample after a DBT with any T2 employed. 3.4 XRD measurements The XRD patterns of top surface layer of the SMAT sample and the coarse-grained sample after a DBT at T1=600℃ for 2h followed by at T2=850℃ for 8h are presented in Fig. 7. For both samples, it is obvious that monophase of Fe2B are detected, which is beneficial for industrial application, owing to its much better toughness when compared to that of FeB phase. The strongest peak can be indexed to 10
the (121) diffraction peak of Fe2B(tetragonal). The intensity of (002) diffraction peak of Fe2B is the second strongest peak(indicated by arrow), which shows that the growth of Fe2B phase exhibits a (002) preferred orientation. Moreover, from the characteristic of diffraction peaks, it can be concluded that the coating phase possesses good crystallinity. Additionally, some weak diffraction peaks can be indexed to the phase of Fe2O3, indicating that the slight oxidation appeared on the sample surface during the tempering process. 3.5 Kinetics study of borided layer growth The growth kinetics of the layer is controlled by the boron diffusion in the Fe2B phase and the growth of borided layer occurs as a consequence of the boron diffusion perpendicular to the specimen surface. According to Eq. (1), the growth of Fe2B layer obeys the parabolic law. The square of borided layer thickness synthesized by a DBT as a function of boronizing time can be described as follows:
d 2 = Kt
(2)
where d is the thickness of borided layer that is measured by the method described in a previous paper [27], K is the growth rate constant depending on boronizing temperature and is calculated from the slope of the d2 versus treatment time graph, which are presented in Fig.8 and Table 1. The relationship between growth rate constant K and boronizing temperature can be expressed by an Arrhenius-type equation as follows [28]: K = K 0 exp(−
Q ) RT
(3)
where K0 is called the collision factor and is a measurement of the effectiveness 11
of collisions between reactive species, Q is an activation factor which indicates the amount of energy(J/mol) required to make the reaction occur. T is the absolute temperature and R is the gas constant(J/(mol⋅K)). Taking the natural logarithm of each side of Eq. (3) gives Eq. (4) as follows: InK = InK 0 −
Q RT
(4)
The plot of InK versus the reciprocal absolute temperature (1/T) is a linear relationship. The activation energy Q and the collision factor K0 can be calculated from the slopes and the intercepts of the graphs in Fig.9, respectively, which are presented in Table 1. The calculated activation energy value for SMAT sample is much lower than that for coarse-grained sample. At the same time, the collision factor K0 of the former is about three orders of magnitude higher than that of the latter, which further confirms the enhancement of the boronizing kinetics of the SMAT sample after a DBT. Moreover, the activation energy for the coarse-grained sample is close to the values obtained by previous reports for AISI H13 steel [8] and AISI P20 steel [29], as showed in Table 2. However, the activation energy for the SMAT sample is approximate to that for AISI 1040 steel which possesses very low alloying elements and it is much easier for this kinds of steel to fabricate borided layers on their surfaces. Based on the results mentioned above, the following empirical equations were derived to estimate the borided layer thicknesses in a given treatment time and temperature.
12
For SMAT sample:
d = 1.6 × 10−2 exp(−
16877.6 )t 1023 ≤ T ≤ 1123 (5) T
−4 For coarse-grained sample: d = 6.4 × 10 exp(−
25181.6 )t 1023 ≤ T ≤ 1123 (6) T
where d is the thickness of borided layer; T is boronizing temperature(K) and t is boronizing time. 3.6 Thermal fatigue property of borided sample From an industrial application point of view, temperature gradients exist on H13 steels from their surfaces to the substrates, which lead to dimensional variation that generates stress and deformation. This gives rise to severe strain in the surface layer of the material, gradually producing thermal fatigue cracks. Fig.10 shows the optical micrographs of surface fatigue cracks of samples without pack boriding treatment after thermal fatigue tests for (a)1000,(b)2000,(c)3000 cycles and pack boriding samples after thermal fatigue tests for (d)1000,(e)2000,(f)3000 cycles. The pack boriding treatment assisted by SMAT was carried out at T1 =600℃ for 2h followed by at T2 =850℃ for 8h. After 1000 cycles, a large amount of long parallel cracks longitudinally distributed on the surface of untreated sample, while only a number of short cracks appeared on the surface of borided sample. When the number of cycles was 2000, along with longitudinal parallel cracks, a lot of short cracks were present transversally on the untreated sample, intersecting with longitudinal parallel cracks. While there were many network-shaped cell structures on the borided sample, which played a role in abruptly propagating the surface cracks [30]. After 3000 cycles, many long and wide cracks existed on the untreated sample. Although the cracks on the 13
borided sample became wide a little when compared to that on the sample after 2000 cycles, the density of them did not increase significantly. Fig.11 displays the SEM images of the cross-section of (a) untreated sample and (b) pack boriding sample assisted by SMAT after thermal fatigue tests for 3000 cycles. For the untreated sample, there were some long and wide cracks that grew into the substrate. On the contrary, for the borided sample, a lot of small and short cracks existed in the coating layer and the interface between coating and substrate, but there were few long cracks which could be detected in the region adjacent to coating layer. Considering the results mentioned above, the pack boriding sample assisted by SMAT possesses much better thermal fatigue resistance than that of the untreated one. It is well known that the oxidation phenomenon can occur in the thermal fatigue tests. During the thermal cycling, the untreated sample can easily be oxidized as it contacted with atmosphere directly. This phenomenon also took place at the internal surface of thermal cracks. The oxide layer with low microhardness could form on the surface layer at 700 , which led to the occurrence of a certain thermal softening [31]. Under the condition of periodical cycles of tensile and compressive stress on the surface, the nucleation and propagation of cracks occurred easily in the region of thermal softening, as shown by our experimental results. As the number of cycles increasing, the thermal cracks could grow and become long and wide, which may expand into substrate. However, the condition of thermal fatigue for the borided sample was different. On the one hand, the oxidation resistance and thermal stability of borided layer are 14
excellent. On the other hand, during heating, the Fe2B phase has lower thermal expansion coefficient(2.9×10-8/ ) than that of substrate(1.35×10-5/ ). Therefore, the expansion of the former was smaller than that of the latter. In order to maintain coating-substrate continuity, the coated substrate could deform less than that of the uncoated one, which could result in a lower stress intensity factor and a lower crack growth rate [32]. In addition, the residual compressive stress in the coatings reduced the stress intensity factor at the nucleated crack tip, which could reduce the crack growth rate in the substrate. Then, only a number of short cracks were observed on the surface of borided sample after 1000 cycles. With the number of cycles increasing, owning to the role of surface oxidation and the difference in thermal expansion coefficient of Fe2B phase with that of the substrate, thermal fatigue cracks could appear in the Fe2B phase. At the same time, the newly formed crack surface which exposed to the atmosphere was subjected to oxidation, due to the oxygen easily penetrated through thermal cracks and pre-existing coating defects [31]. As a result, after 2000 cycles, many network-shaped cell structures were detected on the borided sample. After 3000 cycles, the cracks became wide a little, but most of them were arrested at the region adjacent to the interface of coating and substrate. Obviously, the borided layer could effectively delay the thermal fatigue cracks initiation and impede their propagation. Therefore, the thermal fatigue property of H13 steel with a DBT can be improved remarkably. 4. Conclusions The nanostructured surface layer can be fabricated on H13 steel assisted by SMAT. 15
There are various non-equilibrium GBs in it, along which the diffusion rate of boron is much higher than that along the relaxed GBs in the coarse-grained counterpart. The peak value of boron concentration in the top surface layer for the SMAT sample with pack boriding treatment at 600℃ for 2h is about 6 wt.%, while the peak value for the coarse-grained sample is only about 0.6 wt.%. By means of a DBT at T1 =600℃ for 2h followed by at T2 =850℃ for 8h，the thicknesses of Fe2B phase on the SMAT sample and the coarse-grained counterpart are 40µm and 30µm, respectively. The microhardness of borides formed on the surface of SMAT sample is about 1800HV0.1. In comparison, the hardness of borides is about 1600HV0.1 for the coarse-grained counterpart. The calculated activation energy value for SMAT sample (140.3 kJ/mol) is much lower than that for coarse-grained sample(209.4 kJ/mol). It can be concluded that the boronizing kinetics of the SMAT sample after a DBT at T1=600℃ for 2h followed by at a higher temperature employed can be enhanced remarkably. The thermal fatigue tests show that the borided layer with excellent oxidation resistance and mechanical strength at elevated temperatures could effectively delay the thermal fatigue cracks initiation and impede their propagation. Therefore, the thermal fatigue property of H13 steel with a DBT can be improved remarkably.
Acknowledgements The authors acknowledge the “11th Five” National Science and Technology Support Project of China(Project Number:2007BAE51B04) and Shanghai Leading Academic Discipline Project for finance support(Project Number:S30107). The 16
authors would like to thank P.F.Hu from Instrumental Analysis and Research Center of Shanghai University for the help with the TEM measurements.
References 1. M.B. Karamis, Wear 150 (1991) 331-342. 2. S. Li, R.R. Manory, J. Mater. Sci. 34 (1999)1045-1049. 3. R.L.O. Basso, R.J. Candal, C.A. Figueroa, D. Wisnivesky, F. Alvarez, Surf. Coat. Technol. 203 (2009) 1293-1297. 4. R.D. Torres, C.S.J. Paulo, C. Schmitz, C.J.M. Siqueira, Surf. Coat. Technol. 205 (2010) 1381-1385. 5. S.D. Lu, Z.B. Wang, K. Lu, Mater. Sci. Eng. A 527 (2010) 995-1002. 6. J.C.A. Batista, M.C. Joseph, C. Godoy, A. Matthews, Wear 249 (2002) 971-979. 7. J.A. Davis, P.J. Wilbur, D.L. Williamson, R. Wei, J.J. Vajo, Surf. Coat. Technol. 103–104 (1998) 52-57. 8. K. Genel, Vacuum 80 (2006) 451-457. 9. S. Taktak, Mater. Des. 28 (2007) 1836-1843. 10. W.P. Tong, N.R. Tao, Z.B. Wang, J. Lu, K. Lu, Science 299 (2003) 686-688. 11. W.P. Tong, Z. Han, L.M. Wang, J. Lu, K. Lu, Surf. Coat. Technol. 202 (2008) 4957-4963. 12. Y.M. Lin, J. Lu, L.P. Wang, T. Xu, Q.J. Xue, Acta Mater. 54 (2006) 5599-5605. 13. W.P. Tong, N.R Tao, Z.B Wang, H.W. Zhang, J. Lu, K. Lu, Scr. Mater. 50 (2004) 647-650. 17
14. Z.B. Wang, J. Lu, K. Lu, Acta Mater. 53 (2005) 2081-2089. 15. Z.B. Wang, J. Lu, K. Lu, Surf. Coat. Technol. 201 (2006) 2796-2801. 16. N.R. Tao, Z.B. Wang, W.P. Tong, M.L. Sui, J. Lu, K. Lu, Acta Mater. 50 (2002) 4603-4616. 17. Y. Ivanisenko, W. Lojkowski, R.Z. Valiev, H.J. Fecht, Acta Mater. 51 (2003) 5555-5570. 18. G. Liu, S.C. Wang, X.F. Lou, J. Lu, K. Lu, Scr. Mater. 44 (2001) 1791-1795. 19. Z.B. Wang, N. R.Tao, W. P. Tong, J. Lu, K. Lu, Acta Mater. 51 (2003) 4319-4329. 20. Z.B. Wang, K. Lu, G. Wilde, S. Divinski, Appl. Phys. Lett. 93 (2008) 131904-1-3. 21. Y. Kolobov, G.P. Grabovetskaya, M.B. Ivanov, A.P. Zhilyaev, R.Z.Valiev, Scr. Mater. 44 (2001) 873-878. 22. Y. Amouyal, S.V. Divinski, Y. Estrin, E. Rabkin, Acta Mater. 55 (2007) 59685979. 23. W.D. Callister, Materials Science and Engineering: An Introduction, 5th ed., John Wiley & Sons, New York, 2000, p.106. 24. C.M. Brakman, A.W.J. Gommers, E.J. Mittemeijer, J. Mater. Res. 6 (1989) 13541370. 25. H.C. Fiedler, W.J. Hayes, Metall. Trans. 1 (1970) 1071-1073. 26. B. Chicco, W.E. Borbidge, E. Summerville, Surf. Eng. 14 (1998) 25-30. 27. G. Kartal, S. Timur, V. Sista, O.L. Eryilmaz, A. Erdemir, Surf. Coat. Technol. 206 (2011) 2005-2011. 28. K. Genel, I. Ozbek , C. Bindal, Mater. Sci. Eng. A 347 (2003) 311-314. 18
29. I. Uslu, H. Comert, M. Ipek, F.G. Celebi, O. Ozdemir, C. Bindal, Mater. Des. 28 (2007) 1819-1826. 30. H.Z. Zhi, H. Zhou, L.Q. Ren, X. Tong, H.Y. Shan, L. Liu, Int. J. Fatigue 31 (2009) 468-475. 31. M. Pellizzari, A. Molinari, G. Straffelini, Surf. Coat. Technol. 142–144 (2001) 1109-1115. 32. C.M.D. Starlinga, J.R.T. Branco, Thin Solid Films 308–309 (1997) 436-442.
19
Table 1 Growth rate constant, calculated activation energy and collision factor for test materials as a function of treatment temperatures
Materials
Treatment temperatures T(℃)
Growth rate constant K(m2/s)
SMAT samples
750 800 850 750 800 850
1.33E-14 4.91E-14 5.94E-14 4.15E-15 1.61E-14 3.92E-14
Coarse-grained samples
20
Activation energy Q(kJ/mol) 140.3
Collision factor K0 (m2/s) 2.53E-4
209.4
4.10E-7
Table 2 The comparison of activation energy of the present work and other studies
Materials AISI 1040 steel AISI P20 steel AISI H13 steel AISI H13 steel (SMAT sample) AISI H13 steel (Coarse-grained sample)
Activation energy Q(kJ/mol) 168.0 200.0 186.2 140.3 209.4
21
Reference [29] [29] [8] Present work Present work
Fig.1. Schematic diagram of experimental set up used for thermal fatigue test system: 1 specimen, 2 induction heater, 3 cooling water, 4 numerical control system of heating and cooling, 5 cooling control, 6 high frequency induction generator.
Fig.2. (a) The size and geometry of a thermal fatigue test specimen (unit: mm) and (b) the schematic diagram of two study surfaces.
Fig.3.(a)Typical
bright-field
and
(b)dark-field
TEM
images
showing
the
microstructure of the top surface layer in the SMAT sample and corresponding selected-area electron diffraction pattern (inset) and (c) the statistical distribution of grain size.
Fig.4. The optical micrographs of the cross section of (a)the SMAT sample and (b)the coarse-grained counterpart with pack boriding treatment at 600℃ for 2h; (c) the corresponding boron concentration gradient of the SMAT sample and coarse-grained (CG) counterpart determined by GD-OES tests.
Fig.5. Typical SEM images of the cross section of (a) the coarse-grained sample and (b) the SMAT sample with a DBT at T1 =600℃ for 2h followed by at T2 =850℃ for 8h; (c)the
corresponding
microhardness
gradient
of
the
SMAT
sample
and
coarse-grained(CG) counterpart.
Fig.6. The thicknesses of borided layers formed on the SMAT samples and coarse-grained(CG) counterparts with pack boriding treatment at 600℃ for 2h followed by heating at different higher temperatures for 8h. Fig.7. XRD patterns of (A) the coarse-grained counterpart and (B) the SMAT sample with a DBT at T1 =600℃ for 2h followed by at T2 =850℃ for 8h.
22
Fig.8. Square of borided layers thicknesses of (a) the coarse-grained counterparts and (b) the SMAT samples as a function of boronizing time at different temperatures.
Fig.9. Natural logarithm of growth rate constants of the coarse-grained(CG) counterparts and the SMAT samples as a function of reciprocal of boronizing temperatures.
Fig.10. Optical micrographs of surface fatigue cracks of samples without pack boriding treatment after thermal fatigue tests for (a)1000; (b)2000; (c)3000 cycles and pack boriding samples assisted by SMAT after thermal fatigue tests for (d)1000; (e)2000; (f) 3000 cycles.
Fig.11. Typical SEM images of the cross section of (a) sample without pack boriding treatment and (b) pack boriding sample assisted by SMAT after thermal fatigue tests for 3000 cycles.
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Dear Editors of Journal of Alloys and Compounds: The main highlights of this rearch are as follows: 1. Nanostructured surface layer is fabricated on H13 steel assisted by SMAT. 2. The boronizing kinetics of SMAT sample can be enhanced remarkably.
3. Borided layer can delay fatigue cracks initiation and impede their propagation.
35
S0925-8388(13)03130-7 http://dx.doi.org/10.1016/j.jallcom.2013.12.151 JALCOM 30215
To appear in: Received Date: Revised Date: Accepted Date:
2 July 2013 16 December 2013 16 December 2013
Please cite this article as: H. Yang, X. Wu, Z. Yang, S. Pu, H. Wang, Enhanced boronizing kinetics of alloy steel assisted by surface mechanical attrition treatment, (2013), doi: http://dx.doi.org/10.1016/j.jallcom.2013.12.151
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 proof before it is published in its final 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.
Enhanced boronizing kinetics of alloy steel assisted by surface mechanical attrition treatment Haopeng Yang∗, Xiaochun Wu, Zhe Yang, Shengjun Pu, Hongbin Wang School of Material Science and Engineering, Shanghai University, Shanghai 200072, People’s Republic of China
Abstract: A nanostructured surface layer was fabricated on AISI H13 steel by means of surface mechanical attrition treatment(SMAT). Boronizing behaviors of the SMAT samples were systematically investigated in comparison with their coarse-grained counterparts. The boron diffusion depth of the SMAT sample with pack boriding treatment at 600℃ for 2h was about 8µm, which was much deeper than that of the coarse-grained sample. A much thicker borided layer on the SMAT sample can be synthesized by a duplex boronizing treatment at 600℃ followed by at a higher temperature. The borided layer was composed with monophase of Fe2B and the growth of it exhibited a (002) preferred orientation. Moreover, the activation energy of boron diffusion for the SMAT sample is 140.3kJ/mol, which is much lower than 209.4kJ/mol for the coarse-grained counterpart. The results indicate that the boronizing kinetics can be significantly enhanced in the SMAT sample with a duplex boronizing treatment. Furthermore, the thermal fatigue tests show that the borided layer with excellent oxidation resistance and mechanical strength at elevated temperatures could effectively delay the thermal fatigue cracks initiation and impede their propagation. Therefore, the thermal fatigue property of H13 steel with a duplex boronizing treatment can be improved remarkably. ∗ Corresponding author. Tel.: +86 21 56331153; fax: +86 21 56331461. E-mail addresses: [email protected] (H.P.Yang), [email protected] (X.C.Wu). 1
Keywords: H13 steel; Surface mechanical attrition treatment; Diffusion; Boronizing kinetics 1. Introduction AISI H13 steel is one of the most important hot work die steels. In the service process, it is subjected to high loads and high temperatures, which produce serious damages of tool surfaces because of wear, plastic deformation and thermal fatigue. In order to achieve longer service life, various thermochemical surface hardening treatments have been applied to improve its wear performance, oxidation resistance and thermal fatigue property, such as gas and plasma nitriding [1-3], physical vapor deposition of hard ceramic coating [4], packed powder chromizing treatment [5] and duplex treatments [6]. The technology of duplex surface modification has attracted extensive attention in recent years. Torres et al. [4] used nitriding as pretreatment to prepare TiAlN/TiN coating with different thicknesses on H13 substrate, which could improve coating-substrate adhesion, and the wear rate of the duplex coatings can be decreased significantly. However, duplex treatments usually demand the use of rather complicated expensive apparatus for plasma nitriding and physical or chemical vapor deposition. These disadvantages hinder their wide applications. Boronizing is an effective surface hardening process. The borided layer possesses excellent adhesion to the substrate due to the nature of the thermal diffusion process. The boronizing technique can be carried out in solid, liquid or gaseous medium [7]. Among the various boronizing processes, solid-state pack boriding treatment is the most frequently used. Genel et al. [8] prepared both phases of FeB and Fe2B on H13 2
steel in the solid medium consisting of Ekabor-I powders at 800, 900, 1000℃for periods of 1–5h. Taktak et al. [9] reported the FeB and Fe2B can be formed on H13 steels in slurry salt bath consisting of borax, boric acid and ferrosilicon at temperature range of 800–950℃ for 3, 5 and 7h. However, the disadvantages of pack boriding treatment are the requirements of relatively high processing temperature (850–1000℃) and long process duration to obtain a useful borided layer thickness. Studies have been carried out to improve the efficiency of pack boriding treatment over the past decades. It is well known that borided layer generated in thermochemical treatment depends on boriding condition and on the property of the material itself. Both factors are strongly affected by grain boundaries and defect density. The combination of pack boriding treatment with surface mechanical attrition treatment(SMAT) could provide an alternative approach to observably improve the efficiency of boronizing kinetics. After SMAT, various metals, such as pure Fe [10], 38CrMoAl [11], AISI 321 austenitic stainless [12] and other alloys [13] possess a nanostructured surface layer. Due to the significant enhanced diffusion and chemical reaction kinetics in the formed nanostructured surface layer, the hardened diffusion layer with disperse compounds on the substrate has been fabricated on several ferrous alloys after the subsequent gas nitriding [10,11] or packed powder chromizing treatment [14,15]. In addition, the efficiency of thermochemical treatment has been increased evidently relative to that of the conventional one without SMAT. Here, we report that pack boriding treatment was carried out for the hot work die 3
steel H13 assisted by SMAT. The boronizing kinetics in the nanostructured surface layer fabricated by SMAT was studied. Furthermore, the thermal fatigue property of H13 steel with a duplex boronizing treatment was investigated by means of thermal fatigue tests with 1000, 2000 and 3000 continuous cycles from 20 to 700℃. 2. Experimental procedure 2.1 Sample preparation The chemical compositions of AISI H13 steel used in the experiments contain (wt.%) 0.42C, 4.93Cr, 1.40Mo, 0.98Si, 0.87V, 0.30Mn, 0.018P, 0.005S and balance Fe. Before SMAT, the ferritic steel with the shape(80×80×4mm) was mirror polished. The set-up and procedures of SMAT treatment had been described in a previous paper [16]. In our experiments, the SMAT to H13 steels was performed under vacuum at room temperature for 60min with a vibrating frequency of 50Hz using bearing steel balls with 8mm diameter. Because the sample surface was plastically deformed with high strains and high strain rates, grains in the surface layer are effectively refined. The small samples with dimensions of 15×15×4mm were machined from the bulk sample mentioned above and ultrasonically cleaned in acetone. The pack boriding treatment was as follows: the constituents of the boriding media were B4C(10wt.%), KBF4(5wt.%), charcoal(5wt.%) with the balance of SiC. The SMAT samples and the untreated ones were packed in the powder mix and sealed in a stainless steel container. Boronizing was performed in an electrical resistance furnace at 600℃(T1) for 2h followed by heating at different higher temperatures(T2) for various time. In order to stabilize the nanostructures in the SMAT surface layer, 4
boronizing was carried out at T1, which was useful to form a much thicker borided layer on H13 steel with a subsequent boronizing process at a higher temperature T2. The boronizing temperatures were monitored by a NiCr–NiSi thermocouple with an accuracy of ±2℃. All samples were cooled with furnace. Finally, the borided samples were quenched at 1030℃ and tempered twice at 580℃ in vacuum oven. 2.2 Thermal fatigue tests Thermal fatigue tests were carried out in a high frequency induction test system. The schematic diagram of the thermal fatigue test system is shown in Fig.1. The Uddeholm self-restricted method was used for the thermal fatigue tests. The size and geometry of a thermal fatigue test specimen are shown in Fig.2(a). The integral specimen is composed with two cylinders. The diameter of one cylinder is 5mm and the other is 10mm.Two study areas were prepared symmetrically on the bigger cylinder, as indicated by arrows in Fig.2(b). One study area was the polished surface, the other was the borided surface assisted by SMAT. An induction heating system and a water shower system were used for heating and cooling, respectively. The thermal fatigue tests were carried out from 20 to 700℃.The error of the temperature was about 10℃, considering the fast rates of heating and cooling. The specimen was heated from 20 to 700℃ in 3 seconds and was kept for 1 second to allow the specimen to expand and to reach the equilibrium. Then, the specimen was cooled with water shower for 8 seconds and the temperature could be decreased to about 20℃ rapidly. The total time of one cycle was 12 seconds. The specimen was subjected to repeated compressive and tensile stress. The number of cycles in our experiments 5
included 1000, 2000 and 3000. After thermal fatigue test, the specimen was soaked in HCl solution(10% vol.%) and the oxide skin on specimen can be removed. The crack surfaces were observed with optical microscope to identify crack mode. Moreover, the morphologies of cracks formed in the region beneath the crack surfaces were studied after sectioning the cracked specimen longitudinally. 2.3 Characterization The
X-ray
diffraction(XRD)
measurements
were
obtained
with
a
RigakuD/Max-RBX-ray diffractometer by using CuKa(40kV,40mA) radiation and a secondary beam graphite monochromator. Cross-sectional observations of the borided samples were performed on a JSM-6301F scanning electron microscope(SEM). Microstructural features in the surface layer were characterized by using a field emission transmission electron microscopy(TEM,JEOLJEM-2010F) equipped with an energy dispersive spectral(EDS) analyzer. Thin foil samples for TEM observations were prepared by cutting, grinding and dimpling with a final ion thinning at low temperatures. Microstructures of cross section of the borided layers etched in 4% nital solution were examined by an optical microscope. The hardness distributions of the cross-sectioned borided layers were measured by a Vickers microhardness tester with a load of 100g force. A LECO GDS850A glow-discharge optical emission spectrometer(GD-OES) was used to measure the concentrations of boron element in the surface layers of the borided samples. 3. Results and discussion 3.1 Microstructure of the SMAT sample 6
Surface mechanical attrition is an effective technique to realize the surface self-nanocrystallization on metallic materials. Fig.3(a) and (b) show that the TEM observations in the top surface layer treated by SMAT. Due to the very high strain and strain rate(102 to 103s−1) [16], extremely fine equiaxed ferrite grains with random crystallographic orientations are formed, as revealed by the corresponding selected-area electron diffraction pattern. Carbide phase may be progressively refined into smaller particles or dissolved into the ferrite phase [14,17]. Therefore, no carbides diffraction ring can be detected. Fig.3(c) displays the statistical distribution of grain size. The size of vast majority of grains is at 10-15nm. The grain size increases gradually along the depth in the surface layer. The similar experimental results were also obtained by some previous reports [5,18]. When comparing to the conventional grain boundaries(GBs) in the annealed coarse-grained samples, the GBs formed by the severe plastic deformation are usually in non-equilibrium states which are associated with a high density of defects or a high stored energy [19]. GBs diffusivity in the deformation layer produced by SMAT [20] or equal-channel angular pressing [21,22] will be considerably enhanced relative to that along high angle GBs in the coarse-grained counterpart. It should be noted that the thermal stability of the nanostructured surface layer can be preserved up to 600℃ according to the previous study results [5]. 3.2 Diffusion of boron in the nanostructured surface layer The optical micrographs of the cross section of the SMAT sample and the coarse-grained counterpart with pack boriding treatment at 600℃ for 2h are showed 7
in Fig. 4(a) and (b), respectively. The microstructural morphology in the surface layer with dark contrast region differs from that in the internal substrate for SMAT sample and the boundary between them is distinctive, which indicate that the chemical components and the crystal structures of the two regions are different, resulting in the different corrosion potential for the top surface layer and the internal substrate etched in nital solution. However, the similar result does not obviously exist in the coarse-grained sample. The concentration of boron along the depth from the treated surface of the SMAT sample and the coarse-grained counterpart are presented in Fig. 4(c), which is determined by GD-OES tests. The peak value of concentration in the top surface layer for the SMAT sample is about 6 wt.%, while the peak value for the coarse-grained sample is only about 0.6 wt.%. At the same time, the depth of boron diffusion in SMAT sample is about 8µm, which is much deeper than that in the coarse-grained sample. Based on the experimental results mentioned above, it can be concluded that the thermal diffusion of boron in the nanostructured surface layer fabricated by SMAT can be evidently enhanced when compared to that in the coarse-grained sample with the same experimental parameter. According to the diffusion equation for semi-infinite solid diffusion [23]:
X = α Dmt
(1)
where α is a constant (about 1 in the present case), t is the duration of the pack boriding treatment, and Dm is the mean diffusion coefficient. The measured diffusion depth into the SMAT sample is larger than that into the coarse-grained counterpart, 8
indicating the change of diffusion mechanism in the SMAT surface layer. Various non-equilibrium GBs, along which the diffusion rate of boron is much higher than that along the relaxed GBs in the coarse-grained counterpart, dominate the much enhanced boronizing kinetics in the SMAT surface layer at 600℃. The mean GBs diffusion coefficient calculated from the measured boron diffusion depth in the SMAT sample according to Eq. (1) is 8.9×10-15 m2/s, which is much higher than 1.3×10-15 m2/s for the coarse-grained sample. 3.3 Microstructure and hardness of the borided layer The pack boriding treatment at such a low temperature was found to facilitate stabilizing the nanostructures in the SMAT surface layer, which was useful to form a much thicker borided layer to enhance the surface properties of tool steels with a subsequent boronizing process at a higher temperature. Fig. 5(a) and (b) show the cross-sectional SEM observations for the coarse-grained sample and the SMAT sample after a duplex boronizing treatment(DBT) at T1 =600℃ for 2h followed by at T2 =850℃ for 8h, respectively. It can be seen that the continuous borided layers are on the two sample surfaces. The thicknesses of borided layers on the SMAT sample and the coarse-grained counterpart are 40µm and 30µm, respectively. Neither FeB(orthorhombic) nor Fe2B(tetragonal) has cubic crystal symmetry [24]. Although H13 steel possesses a certain amount of alloying elements, boron diffusion in the borided layer is of anisotropic nature at high temperatures. Therefore, the morphologies of coating layers exhibit tooth-like growth into substrate. The microhardness gradients for the two samples are showed in Fig. 5(c). For both samples, the graphs include three distinct regions, which are borided layer, transition zone and matrix that is not affected by boron diffusion. The microhardness of borides formed on the surface of SMAT sample is about 1800HV0.1 and it is roughly 3.5 times harder than that of the substrate. In comparison, the microhardness of borides is about 1600HV0.1 for the coarse-grained counterpart. Owing to significant refined grain sizes 9
in the nanostructured surface layer of SMAT sample, the refined grains of borides on H13 steel could be synthesized with a DBT. However, after the borided samples were quenched at 1030℃ and tempered twice at 580℃, the refined grains of borides on SMAT sample could grow into larger crystalline grains, and the difference between the grain sizes of borides on SMAT sample and that on coarse-grained counterpart is not significantly obvious. At the same time, the continuous and dense microscopic structures of borides on them are not evidently different. Therefore, there is not a substantial difference in the microhardness gradients between the SMAT and coarse-grained samples with a DBT. In addition, the microhardness of transition zone is a little softer than that of the substrate. Similar results were reported for H11, H12 and H13 steels [25,26]. The decrease in hardness can be ascribed to concentration of the silicon atom in front of borided layer, as silicon is not soluble in it. It is well known that silicon is a ferrite-forming element and the high concentration of silicon ahead of the borided layer suggests that this region remained in ferrite at the austenitizing temperature and consequently did not subsequently transform into martensite as did in the internal substrate during the quenching process [25]. As shown in Fig. 6, based on the pack boriding treatment at T1=600℃ for 2h, a much thicker borided layer was formed on the SMAT sample than that on the coarse-grained counterpart after the boronizing treatment at any T2 from 750-850℃ for 8h. The thicknesses in both samples increase with an increasing T2. It can be seen that the low temperature treatment at 600℃ will enhance the boronizing kinetics of the SMAT sample after a DBT with any T2 employed. 3.4 XRD measurements The XRD patterns of top surface layer of the SMAT sample and the coarse-grained sample after a DBT at T1=600℃ for 2h followed by at T2=850℃ for 8h are presented in Fig. 7. For both samples, it is obvious that monophase of Fe2B are detected, which is beneficial for industrial application, owing to its much better toughness when compared to that of FeB phase. The strongest peak can be indexed to 10
the (121) diffraction peak of Fe2B(tetragonal). The intensity of (002) diffraction peak of Fe2B is the second strongest peak(indicated by arrow), which shows that the growth of Fe2B phase exhibits a (002) preferred orientation. Moreover, from the characteristic of diffraction peaks, it can be concluded that the coating phase possesses good crystallinity. Additionally, some weak diffraction peaks can be indexed to the phase of Fe2O3, indicating that the slight oxidation appeared on the sample surface during the tempering process. 3.5 Kinetics study of borided layer growth The growth kinetics of the layer is controlled by the boron diffusion in the Fe2B phase and the growth of borided layer occurs as a consequence of the boron diffusion perpendicular to the specimen surface. According to Eq. (1), the growth of Fe2B layer obeys the parabolic law. The square of borided layer thickness synthesized by a DBT as a function of boronizing time can be described as follows:
d 2 = Kt
(2)
where d is the thickness of borided layer that is measured by the method described in a previous paper [27], K is the growth rate constant depending on boronizing temperature and is calculated from the slope of the d2 versus treatment time graph, which are presented in Fig.8 and Table 1. The relationship between growth rate constant K and boronizing temperature can be expressed by an Arrhenius-type equation as follows [28]: K = K 0 exp(−
Q ) RT
(3)
where K0 is called the collision factor and is a measurement of the effectiveness 11
of collisions between reactive species, Q is an activation factor which indicates the amount of energy(J/mol) required to make the reaction occur. T is the absolute temperature and R is the gas constant(J/(mol⋅K)). Taking the natural logarithm of each side of Eq. (3) gives Eq. (4) as follows: InK = InK 0 −
Q RT
(4)
The plot of InK versus the reciprocal absolute temperature (1/T) is a linear relationship. The activation energy Q and the collision factor K0 can be calculated from the slopes and the intercepts of the graphs in Fig.9, respectively, which are presented in Table 1. The calculated activation energy value for SMAT sample is much lower than that for coarse-grained sample. At the same time, the collision factor K0 of the former is about three orders of magnitude higher than that of the latter, which further confirms the enhancement of the boronizing kinetics of the SMAT sample after a DBT. Moreover, the activation energy for the coarse-grained sample is close to the values obtained by previous reports for AISI H13 steel [8] and AISI P20 steel [29], as showed in Table 2. However, the activation energy for the SMAT sample is approximate to that for AISI 1040 steel which possesses very low alloying elements and it is much easier for this kinds of steel to fabricate borided layers on their surfaces. Based on the results mentioned above, the following empirical equations were derived to estimate the borided layer thicknesses in a given treatment time and temperature.
12
For SMAT sample:
d = 1.6 × 10−2 exp(−
16877.6 )t 1023 ≤ T ≤ 1123 (5) T
−4 For coarse-grained sample: d = 6.4 × 10 exp(−
25181.6 )t 1023 ≤ T ≤ 1123 (6) T
where d is the thickness of borided layer; T is boronizing temperature(K) and t is boronizing time. 3.6 Thermal fatigue property of borided sample From an industrial application point of view, temperature gradients exist on H13 steels from their surfaces to the substrates, which lead to dimensional variation that generates stress and deformation. This gives rise to severe strain in the surface layer of the material, gradually producing thermal fatigue cracks. Fig.10 shows the optical micrographs of surface fatigue cracks of samples without pack boriding treatment after thermal fatigue tests for (a)1000,(b)2000,(c)3000 cycles and pack boriding samples after thermal fatigue tests for (d)1000,(e)2000,(f)3000 cycles. The pack boriding treatment assisted by SMAT was carried out at T1 =600℃ for 2h followed by at T2 =850℃ for 8h. After 1000 cycles, a large amount of long parallel cracks longitudinally distributed on the surface of untreated sample, while only a number of short cracks appeared on the surface of borided sample. When the number of cycles was 2000, along with longitudinal parallel cracks, a lot of short cracks were present transversally on the untreated sample, intersecting with longitudinal parallel cracks. While there were many network-shaped cell structures on the borided sample, which played a role in abruptly propagating the surface cracks [30]. After 3000 cycles, many long and wide cracks existed on the untreated sample. Although the cracks on the 13
borided sample became wide a little when compared to that on the sample after 2000 cycles, the density of them did not increase significantly. Fig.11 displays the SEM images of the cross-section of (a) untreated sample and (b) pack boriding sample assisted by SMAT after thermal fatigue tests for 3000 cycles. For the untreated sample, there were some long and wide cracks that grew into the substrate. On the contrary, for the borided sample, a lot of small and short cracks existed in the coating layer and the interface between coating and substrate, but there were few long cracks which could be detected in the region adjacent to coating layer. Considering the results mentioned above, the pack boriding sample assisted by SMAT possesses much better thermal fatigue resistance than that of the untreated one. It is well known that the oxidation phenomenon can occur in the thermal fatigue tests. During the thermal cycling, the untreated sample can easily be oxidized as it contacted with atmosphere directly. This phenomenon also took place at the internal surface of thermal cracks. The oxide layer with low microhardness could form on the surface layer at 700 , which led to the occurrence of a certain thermal softening [31]. Under the condition of periodical cycles of tensile and compressive stress on the surface, the nucleation and propagation of cracks occurred easily in the region of thermal softening, as shown by our experimental results. As the number of cycles increasing, the thermal cracks could grow and become long and wide, which may expand into substrate. However, the condition of thermal fatigue for the borided sample was different. On the one hand, the oxidation resistance and thermal stability of borided layer are 14
excellent. On the other hand, during heating, the Fe2B phase has lower thermal expansion coefficient(2.9×10-8/ ) than that of substrate(1.35×10-5/ ). Therefore, the expansion of the former was smaller than that of the latter. In order to maintain coating-substrate continuity, the coated substrate could deform less than that of the uncoated one, which could result in a lower stress intensity factor and a lower crack growth rate [32]. In addition, the residual compressive stress in the coatings reduced the stress intensity factor at the nucleated crack tip, which could reduce the crack growth rate in the substrate. Then, only a number of short cracks were observed on the surface of borided sample after 1000 cycles. With the number of cycles increasing, owning to the role of surface oxidation and the difference in thermal expansion coefficient of Fe2B phase with that of the substrate, thermal fatigue cracks could appear in the Fe2B phase. At the same time, the newly formed crack surface which exposed to the atmosphere was subjected to oxidation, due to the oxygen easily penetrated through thermal cracks and pre-existing coating defects [31]. As a result, after 2000 cycles, many network-shaped cell structures were detected on the borided sample. After 3000 cycles, the cracks became wide a little, but most of them were arrested at the region adjacent to the interface of coating and substrate. Obviously, the borided layer could effectively delay the thermal fatigue cracks initiation and impede their propagation. Therefore, the thermal fatigue property of H13 steel with a DBT can be improved remarkably. 4. Conclusions The nanostructured surface layer can be fabricated on H13 steel assisted by SMAT. 15
There are various non-equilibrium GBs in it, along which the diffusion rate of boron is much higher than that along the relaxed GBs in the coarse-grained counterpart. The peak value of boron concentration in the top surface layer for the SMAT sample with pack boriding treatment at 600℃ for 2h is about 6 wt.%, while the peak value for the coarse-grained sample is only about 0.6 wt.%. By means of a DBT at T1 =600℃ for 2h followed by at T2 =850℃ for 8h，the thicknesses of Fe2B phase on the SMAT sample and the coarse-grained counterpart are 40µm and 30µm, respectively. The microhardness of borides formed on the surface of SMAT sample is about 1800HV0.1. In comparison, the hardness of borides is about 1600HV0.1 for the coarse-grained counterpart. The calculated activation energy value for SMAT sample (140.3 kJ/mol) is much lower than that for coarse-grained sample(209.4 kJ/mol). It can be concluded that the boronizing kinetics of the SMAT sample after a DBT at T1=600℃ for 2h followed by at a higher temperature employed can be enhanced remarkably. The thermal fatigue tests show that the borided layer with excellent oxidation resistance and mechanical strength at elevated temperatures could effectively delay the thermal fatigue cracks initiation and impede their propagation. Therefore, the thermal fatigue property of H13 steel with a DBT can be improved remarkably.
Acknowledgements The authors acknowledge the “11th Five” National Science and Technology Support Project of China(Project Number:2007BAE51B04) and Shanghai Leading Academic Discipline Project for finance support(Project Number:S30107). The 16
authors would like to thank P.F.Hu from Instrumental Analysis and Research Center of Shanghai University for the help with the TEM measurements.
References 1. M.B. Karamis, Wear 150 (1991) 331-342. 2. S. Li, R.R. Manory, J. Mater. Sci. 34 (1999)1045-1049. 3. R.L.O. Basso, R.J. Candal, C.A. Figueroa, D. Wisnivesky, F. Alvarez, Surf. Coat. Technol. 203 (2009) 1293-1297. 4. R.D. Torres, C.S.J. Paulo, C. Schmitz, C.J.M. Siqueira, Surf. Coat. Technol. 205 (2010) 1381-1385. 5. S.D. Lu, Z.B. Wang, K. Lu, Mater. Sci. Eng. A 527 (2010) 995-1002. 6. J.C.A. Batista, M.C. Joseph, C. Godoy, A. Matthews, Wear 249 (2002) 971-979. 7. J.A. Davis, P.J. Wilbur, D.L. Williamson, R. Wei, J.J. Vajo, Surf. Coat. Technol. 103–104 (1998) 52-57. 8. K. Genel, Vacuum 80 (2006) 451-457. 9. S. Taktak, Mater. Des. 28 (2007) 1836-1843. 10. W.P. Tong, N.R. Tao, Z.B. Wang, J. Lu, K. Lu, Science 299 (2003) 686-688. 11. W.P. Tong, Z. Han, L.M. Wang, J. Lu, K. Lu, Surf. Coat. Technol. 202 (2008) 4957-4963. 12. Y.M. Lin, J. Lu, L.P. Wang, T. Xu, Q.J. Xue, Acta Mater. 54 (2006) 5599-5605. 13. W.P. Tong, N.R Tao, Z.B Wang, H.W. Zhang, J. Lu, K. Lu, Scr. Mater. 50 (2004) 647-650. 17
14. Z.B. Wang, J. Lu, K. Lu, Acta Mater. 53 (2005) 2081-2089. 15. Z.B. Wang, J. Lu, K. Lu, Surf. Coat. Technol. 201 (2006) 2796-2801. 16. N.R. Tao, Z.B. Wang, W.P. Tong, M.L. Sui, J. Lu, K. Lu, Acta Mater. 50 (2002) 4603-4616. 17. Y. Ivanisenko, W. Lojkowski, R.Z. Valiev, H.J. Fecht, Acta Mater. 51 (2003) 5555-5570. 18. G. Liu, S.C. Wang, X.F. Lou, J. Lu, K. Lu, Scr. Mater. 44 (2001) 1791-1795. 19. Z.B. Wang, N. R.Tao, W. P. Tong, J. Lu, K. Lu, Acta Mater. 51 (2003) 4319-4329. 20. Z.B. Wang, K. Lu, G. Wilde, S. Divinski, Appl. Phys. Lett. 93 (2008) 131904-1-3. 21. Y. Kolobov, G.P. Grabovetskaya, M.B. Ivanov, A.P. Zhilyaev, R.Z.Valiev, Scr. Mater. 44 (2001) 873-878. 22. Y. Amouyal, S.V. Divinski, Y. Estrin, E. Rabkin, Acta Mater. 55 (2007) 59685979. 23. W.D. Callister, Materials Science and Engineering: An Introduction, 5th ed., John Wiley & Sons, New York, 2000, p.106. 24. C.M. Brakman, A.W.J. Gommers, E.J. Mittemeijer, J. Mater. Res. 6 (1989) 13541370. 25. H.C. Fiedler, W.J. Hayes, Metall. Trans. 1 (1970) 1071-1073. 26. B. Chicco, W.E. Borbidge, E. Summerville, Surf. Eng. 14 (1998) 25-30. 27. G. Kartal, S. Timur, V. Sista, O.L. Eryilmaz, A. Erdemir, Surf. Coat. Technol. 206 (2011) 2005-2011. 28. K. Genel, I. Ozbek , C. Bindal, Mater. Sci. Eng. A 347 (2003) 311-314. 18
29. I. Uslu, H. Comert, M. Ipek, F.G. Celebi, O. Ozdemir, C. Bindal, Mater. Des. 28 (2007) 1819-1826. 30. H.Z. Zhi, H. Zhou, L.Q. Ren, X. Tong, H.Y. Shan, L. Liu, Int. J. Fatigue 31 (2009) 468-475. 31. M. Pellizzari, A. Molinari, G. Straffelini, Surf. Coat. Technol. 142–144 (2001) 1109-1115. 32. C.M.D. Starlinga, J.R.T. Branco, Thin Solid Films 308–309 (1997) 436-442.
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Table 1 Growth rate constant, calculated activation energy and collision factor for test materials as a function of treatment temperatures
Materials
Treatment temperatures T(℃)
Growth rate constant K(m2/s)
SMAT samples
750 800 850 750 800 850
1.33E-14 4.91E-14 5.94E-14 4.15E-15 1.61E-14 3.92E-14
Coarse-grained samples
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Activation energy Q(kJ/mol) 140.3
Collision factor K0 (m2/s) 2.53E-4
209.4
4.10E-7
Table 2 The comparison of activation energy of the present work and other studies
Materials AISI 1040 steel AISI P20 steel AISI H13 steel AISI H13 steel (SMAT sample) AISI H13 steel (Coarse-grained sample)
Activation energy Q(kJ/mol) 168.0 200.0 186.2 140.3 209.4
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Reference [29] [29] [8] Present work Present work
Fig.1. Schematic diagram of experimental set up used for thermal fatigue test system: 1 specimen, 2 induction heater, 3 cooling water, 4 numerical control system of heating and cooling, 5 cooling control, 6 high frequency induction generator.
Fig.2. (a) The size and geometry of a thermal fatigue test specimen (unit: mm) and (b) the schematic diagram of two study surfaces.
Fig.3.(a)Typical
bright-field
and
(b)dark-field
TEM
images
showing
the
microstructure of the top surface layer in the SMAT sample and corresponding selected-area electron diffraction pattern (inset) and (c) the statistical distribution of grain size.
Fig.4. The optical micrographs of the cross section of (a)the SMAT sample and (b)the coarse-grained counterpart with pack boriding treatment at 600℃ for 2h; (c) the corresponding boron concentration gradient of the SMAT sample and coarse-grained (CG) counterpart determined by GD-OES tests.
Fig.5. Typical SEM images of the cross section of (a) the coarse-grained sample and (b) the SMAT sample with a DBT at T1 =600℃ for 2h followed by at T2 =850℃ for 8h; (c)the
corresponding
microhardness
gradient
of
the
SMAT
sample
and
coarse-grained(CG) counterpart.
Fig.6. The thicknesses of borided layers formed on the SMAT samples and coarse-grained(CG) counterparts with pack boriding treatment at 600℃ for 2h followed by heating at different higher temperatures for 8h. Fig.7. XRD patterns of (A) the coarse-grained counterpart and (B) the SMAT sample with a DBT at T1 =600℃ for 2h followed by at T2 =850℃ for 8h.
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Fig.8. Square of borided layers thicknesses of (a) the coarse-grained counterparts and (b) the SMAT samples as a function of boronizing time at different temperatures.
Fig.9. Natural logarithm of growth rate constants of the coarse-grained(CG) counterparts and the SMAT samples as a function of reciprocal of boronizing temperatures.
Fig.10. Optical micrographs of surface fatigue cracks of samples without pack boriding treatment after thermal fatigue tests for (a)1000; (b)2000; (c)3000 cycles and pack boriding samples assisted by SMAT after thermal fatigue tests for (d)1000; (e)2000; (f) 3000 cycles.
Fig.11. Typical SEM images of the cross section of (a) sample without pack boriding treatment and (b) pack boriding sample assisted by SMAT after thermal fatigue tests for 3000 cycles.
23
Fig.1
24
Fig.2
25
Fig.3
26
Fig.4
27
Fig.5
28
Fig.6
29
Fig.7
30
Fig.8
31
Fig.9
32
Fig.10
33
Fig.11
34
Dear Editors of Journal of Alloys and Compounds: The main highlights of this rearch are as follows: 1. Nanostructured surface layer is fabricated on H13 steel assisted by SMAT. 2. The boronizing kinetics of SMAT sample can be enhanced remarkably.
3. Borided layer can delay fatigue cracks initiation and impede their propagation.
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