铬-钼- V钢的回火脆性 外文翻译.doc

铬-钼- V钢的回火脆性可逆回火脆性（RTB） 在500-650温度范围内回火或缓慢冷却的钢的脆性，它被认为是造成在前奥氏体晶粒边界形成杂质（P，锑，锡，砷） 1-4的原因。然而，有资料显示5，不仅有这些进程，而且有其他进程，像在500-600淬火钢有助于回火脆性在发生。 在这项工作的关注是淬火铬钼合金，以防止在钒和磷的含量RTB的钢回火脆化。 热件1（0.35的V，0.015P）重四十一吨是伪造的，以一个550毫米的大小, 15毫米厚的板材被切断他们。另外，在一个100公斤的感应炉加热熔化。 钒被添加到在一个重达16公斤的铸块中，它被锻造、轧制成10毫米厚的金属板。在所有的10-15毫米厚加热板放在油中980 （1小时）淬火。从板淬火和在100-760回火10小时后制备样品。 经过淬火和低温回火，所有加热件呈现一个典型的马氏体结构，但在高（超过600摄氏度）的温度回火后得到回火索氏体结构。在实验室加热的奥氏体晶粒尺寸较小（8-9级）比商业热（4级），它们的结构部件具有更大的分散性和均匀性的。 脆性是从Tso 和 Ttemper变化中来确定， Ttemper是淬火钢的回火温度，Tso是韧-脆转变温度，它具有最完整的脆化特征。Tso的确定是通过5527.5毫米的切口1毫米深（根半径0.25毫米）冲击试验样品。Tso被认为是在断裂50纤维下的测试温度。拉伸强度通过5个直径为3毫米的样品确定的为20。 图1显示了热件1回火温度与力学性能的变化。这种热铸块的力学性能在100-600 回火后几乎不变。当回火温度从600-760度提高时，强度特性急剧下降，而韧性增加。 更改影响回火钢的稳定的情况下钒浓度，因为不含钒的钢的强度会下降至约500 。 热件1开始时回火温度大约300，Tso的增加达到500-600度的最高值。如同淬火条件对峰值相比增加了100 。回火温度进一步提高，Tso 会下降，这与削弱开始相吻合。 对于热件2-7的Tso和 Ttemper变化的总体特点是相同的，尽管在较低温达到峰巅。这类钢在是Tso90条件冷却，低于热1大概 - 70 10 。在730 后的Tso锻炼价值与热2-7（-110至-130 ）实际上是相同。 它的峰高和立场取决于钒钢的内容。当钒浓度变化从0到0.55的高峰上升60 ，同时转移100 。比较热件1和3，6，在史前的冶金学上有区别的，但相同钒含量（0.3）是相同的，人们可以看到，在Tso峰地区的增长几乎是与淬火条件相同。 在0.005和0.022的磷的热处理后的测试中，磷对Tso没有什么影响。 在回火期间的钢的力学性能改变显然取决于精细结构的变化。图3显示了在热淬火Tso显微条件后，在600回火，对应于对强度特性极值边缘与Tso高峰在760 的回火之后，强度和Tso达到最低的。 淬火后的结构由带有位错的板条马氏体组成。在板条彼此间略有错误导向，平均宽度0.3和长度5，并归为5 5。板条充满均匀分布的密度位错011cm-2。硬质合金阶段没有观察到淬火钢。 在600回火10小时的后混乱阵列的整体性质和晶体的碎片将被保留。木板条的平均大小（宽度和长度）它的平均规模保持不变。唯一的变化是明显的沉淀分散的碳化物阶段。粒子的大小是150至200 A，平均密度为1015cm-2。这些沉淀在混乱通过板条马氏体空间分布。带长条形直径为250A和2103A的较大沉淀物位于深的晶体的边界附近（板条马氏体）。这种类型的技术被作为MTCa电石鉴定。 在760 回火后急剧变化的钢的结构：位错是不规则的。大部分分布在整提失调，改建后更适合了，与位错领域相吻合。单元格的平均规模为0.5。该阶段碳化物形态的也有变化。微细分散的大小圆形的沉淀250 A的位错在位于网络的连接处。随着他们有更大的圆形沉淀2103 A位于边界的十字路口。 析出的碳化物回火后的存在也证实了热物理化学相分析热件1。当回火温度由原来300-600度的渗碳体碳化物的数量逐渐减少， M7C3增加。强有力的碳化物含量形成的残留增加元素：铬，钼，和钒.，这个过程在500-600度是特别活跃。这很可能是VC和 MO2C膜类型不理化分析检测，在高度分散的碳化物阶段也有更多的热力学稳定相6。 因此，氢脆回火过程中铬钼钒钢，作为对Tso与Ttemper出现明显的高峰表现，从我们的实验结果。它的峰值并没有关联的磷在钢存在，但与钒的浓度各不相同，在碳化物的形成，与继承发展脱位阵列的温度范围内位置。Tso达到高峰时的温度在Fe3C的变换更稳，与M7C3类似的现象已在铬钢观察5。脆化可能是由于体积和边界的影响。在铬钼钒钢的的情况二次硬化，体现在在 Ttemper = 300 - 500 性能的增强可能与这些相关的影响有关。据了解，7，随着晶粒尺寸不变（或板条马氏体晶粒常量大小）Tso通过Tso=y+ C与低屈服强度y（0.2相关的关系）同期相联系，其中C和为常数。热件1Tso呈线性关系中的屈服强度为o.2大的变化范围，并只有在400-600度，特殊的碳化物开始制定改变，有偏离线性的关系。当然，这些过程首先影响边界地区，那里的条件是为硬质合金阶段准备，由于晶体结构的几何缺陷有利合金相的形成。可以设想，Tso峰值是由于在边界条件的变化。这个假设是证实了电子显微镜分析结果：在矩阵（位错密度，碎片大小条件）引起材料的8，9实力的增强，保持不变直至Ttemper = 600 。这也证实了结构进行检查和对回火温度与屈服强度的变化。因此，如果发生任何更改的边界，增加脆化，那么，Tso同Ttemper变化将有虚线的形式。所有解释RTB1，2，10的理论都是建立在对边界的影响起主导作用基础上的。然而，前奥氏体晶粒边界损害并非我们的典型调查（热件1纹裂在淬火、回火后显现出来），没有磷的影响，这显然是对钼的存在钢铁中的解释。由此可见，由实验数据CMoV钢回火过程中的脆性主要受碳化物形成的影响，更加准确地重建碳化物和的这些过程影响因素。这原种影响的理可以概括介绍如下，随回火温度的增加渗碳体开始凝聚，在250-350开始沉淀。更有效地遏制电石，在钢中容纳更多的碳化物成形元素比铁中350左右开始的特殊碳化物的原子核钢的成形，尤其是M7C3 8。沉淀物均匀分布在整个体积的位错位置，低角度的板条马氏体的界限，高角度的板条殖民地边界，他们加强了矩阵和削弱（脆化）边界。 在二次硬化温度时，很可能最大程度的削弱碳化物与基体或与沉淀、连贯的最大密度的边界，i.e.。如元素钒，促进碳化物细化，从而增加二次硬化8，增加脆化。 在沉淀和聚结硬质合金阶段，出现了混乱的位错的同时，边界变得更加完善和矩阵被削弱。这两种效应导致韧性、脆性转变温度下降，这是从实际观察的结构。结论1。 15Kh3MFA类型的钢都容易脆化，在给定的回火温度下达到高峰。在高峰期（Tso）的上限，在正值温度下材料开始削弱。 2。该峰的高度和回火温度也相应增加时，几乎线性钒浓度从0提高到0.55，但他们是在0.005-0.022的磷浓度限制下表现出来的。 3。回火钢的脆性取决于在碳化物相变（从渗碳特殊碳化物），保留的位错优先发生在碎片的边界。文献引用1.L. M. Utevskii, Temper Brittleness of Steels in Russian, Metallurgizdat, Moscow (1961),p. 138.2. P. B. Mikhailov-Mikheev, Thermal Embrittlement of Steels in Russian, Mashgiz, Moscow-Leningrad (1956), p. 56.3. J. Hollomon, Trans. ASM, 36, 473 (1946).4. E. Houdremont, Special Steels Russian translation, Vol. I, Metallurgiya, Moscow (1966),p. 455.5. V. A. Korablev, Yu. I. Ustinovshchikov, and I. G. Khatskelevich, Embrittlement of chromium steels with formation of special carbides, Metalloved. Term. Obrab. Met., No. I,16 (1975).6. A. P. Gulyaev, I. K. Kupalova, and V. A. Landa, Method and results of phase analysisof hlgh-speed steels, Zavod. Lab., No. 3, 298 (1965).7. J. Heslop and N. Perch, Phil. Mag., , No. 34, 1128 (1958).8. V. V. Rybin et al., The mechanism of hardening of sorbite-hardening steel and the possibility of determining it theoretically and experimentally, in: Metal Science inRussian, No. 17, Sudostroenie, Leningrad (1973), p. 105.9L. K. Gordienko, Substrucutral Hardening of Metals and Alloys in Russian, Nauka, Moscow(1973), p. 64.10E. E. Glikman et al., Nature of reversible temper brittleness, Fiz. Met. Metalloved.,36, 365 (1973).11. Oliver WC, Pharr GM (2004) J Mater Res 19:312. Kim JY, Lee BW, Read DT, Kwon D (2005) Scr Mater 52:35313. Kim JY, Lee JS, Lee KW, Kim KH, Kwon D (2006) Key Eng Mater 326328:48714. Kim JY, Lee JJ, Lee YH, Jang JI, Kwon D (2006) J Mater Res 21:297515. Kim JY, Kang SK, Lee JJ, Jang JI, Lee YH, Kwon D (2007) Acta Mater 55:355516. Dowling NE (1993) Mechanical behavior of materials. Prentice Hall, Englewood Cliffs17. Kim JY, Lee KW, Lee JS, Kwon D (2006) Surf Coat Technol 201:427818. DIN 17175-79 (1979) Seamless steel tubes for elevated temperatures19. Ahn JH, Kwon D (2001) J Mater Res 16:317020. Dieter GE (1988) Mechanical metallurgy. McGraw-Hill, Singapore外文原文TEMPER BRITTLENESS OF Cr-Mo- V STEELReversible temper brittleness (RTB) embrittlement of steels during tempering or slow cooling in the temperature range of 500-650 is considered to result from the formation of impurity segregates (P, Sb, Sn, As) in prior austenite grain boundaries 1-4. However, there are data indicating 5 that not only these processes but other processes favoring embrittlement occur in quenched steel during tempering at 500-600 .This work concerns embrittlement during tempering of quenched chromium steels alloyed with molybdenum to prevent RTB in relation to the vanadium and phosphorus concentrations.Ingots of heat 1 (0.35% V, 0.015% P) weighing 41 tons were forged to a size of 550 mm. Plates 15 mm thick were cut from them. The other heats were melted in a 100-kg induction furnace with use of ZhS-0 iron in the charge.Vanadium was added to the steel during pouring of an ingot weighing 16 kg, which was forged and rolled to a plate 10 mm thick. All heats in the form of plates 10-15 mm thick were oil quenched from 980 (1 h). Samples were prepared from the plates after quenching and after tempering at 100-760 for 10 h.After quenching and low-temperature tempering, all heats had a typical martensitic structure, but after high-temperature tempering (above 600 ) a sorbite structure. The austenite grain size of the laboratory heats was smaller (grade 8-9) than in the commercial heat (grade 4), with greater dispersity and homogeneity of the structural components.Embrittlement was determined from the variation of Tso with Ttemper , where Ttemper is the tempering temperature of the quenched steel and Tso is the ductile- brittle transition temperature, which most completely characterizes embrittlement.Tso was determined on impact test samples 5 5 27.5 mm with a notch 1 mm deep (root radius 0.25 mm). Tso was taken as the testing temperature at which the fracture was 50% fibrous. The tensile strength was determined at 20 on five samples 3 mm in diameter.Figure 1 shows the variation of the mechanical properties of heat 1 with the tempering temperature. The mechanical properties of this heat are almost constant after tempering at 100-600 . When the tempering temperature is raised from 600-760 the strength characteristics decrease sharply, while the ductile characteristics increase.Changing the vanadium concentration affects the stability of the steel during tempering. For the steel without vanadium the strength begins to decrease around 500 . For heat 1, beginning with tempering at temperatures around 300 , Tso increases and reaches a maximum value at 500-600 . As compared with the quenched condition the increase of Tso at the peak is 100 . With further increase of the tempering temperature Tso decreases, which coincides with the beginning of weakening.For heats 2-7 the overall character of the variation of Tso with Ttemper is the same, although the peaks occur at lower temperatures. For the steels in the quenched condition Tso is 90 lower than for heat 1, and amounts to -70 10 . After tempering at 730 the value of Tso is practically the same for heats 2-7 (-110 to -130).The height of the peak and its position depends on the vanadium content of the steel. When the vanadium concentration is changed from 0 to 0.55% the peak rises %60 and at the same time shifts 100 . Comparing heats 1 and 3, 6, differing in their metallurgical prehistory but similar in vanadium content (0.3%), one can see that the increase of Tso in the region of the peak is almost the same as for the quenched condition.In amounts of 0.005 and 0.022%, phosphorus has no effect on Tso after the heat treatments tested.The changes in the mechanical properties of the steel during tempering evidently depend on changes in fine structure. Figure 3 shows the microstructure of heat in the quenched condition, after tempering at 600 , corresponding to the edge of the plateau of the strength characteristics and the peak of Tso, and after tempering at 760 , where the strength and Tso are lowest.After quenching, the structure consists of lath martensite with well-developed dislocation arrays. The laths are slightly misoriented with respect to each other, with an average width of 0.3 and length 5, and are grouped in colonies 5 5. The laths are filled with evenly distributed dislocations with a density 1011cm-2. No carbide phase was observed in the quenched steel.After tempering at 600 for 10 h the overall character of the dislocation arrays and the fragmentation of the crystals are retained. The average size of the laths (width and length)and the average size of the colonies remain unchanged. The only noticeable change is the precipitation of finely dispersed carbide phases. The size of the particles is 150-200 A and the average density 1015cm-2. They are precipitated on dislocations and evently distributed through the bulk of the martensite laths. Larger precipitates with the shape of needles (platelets) 250 A in diameter and 2103A long are located near highly misoriented boundaries (grains, colonies of martensite laths, the most misoriented laths, etc.) and in the boundaries themselves. Precipitates of this type were identified by microdiffraction techniques as MTCa carbide. The structure of the steel changes sharply after tempering at 760 : the dislocations are polygonized. The dislocations distributed throughout the bulk are rebuilt into energetically more suitable configurationscell wallsnot coinciding with long-range fields. The average size of the cells is 0.5 . The morphology of the carbide phases also changes. Finely dispersed rounded precipitates with a size of 250 A are located in the junctions of the dislocation network. Along with them there are larger rounded precipitates2103 located at the intersections of cell boundaries.The presence of carbide precipitates after tempering was also confirmed by physicochemical phase analysis of heat 1. When the tempering temperature is raised from 300-600 the quantity of Fe3C carbide gradually decreases and the quantity of M7C3 increases. The concentration of strong carbide-forming elements in the residues increases: Cr, Mo, and V. This process is particularly well developed at 500-600 . It is highly probable that among the highly dispersed carbide phases there are also more thermodynamically stable phases of the Mo2C and VC type not detected by physicochemical analysis 6. It follows from our experimental results that embrittlement occurs during tempering of CrMoV steels, manifest as a distinct peak on the curve of Tso vs Ttemper. The peak is not associated with the presence of phosphorus in the steel but varies with the concentration of vanadium and is located in the temperature range of carbide formation, which develops with inheritance of the dislocation arrays. Tso reaches a peak at those temperatures where Fe3C transforms to more stable M7C3. A similar phenomenon has been observed in chromium steels 5. Embrittlement may be due to both volume and boundary effects. In the case of CrMoV steels the secondary hardening manifest in the increase of the strength at Ttemper = 300500 may be associated with the first of these effects. It is known 7 that with a constant grain size (or constant size of colonies of martensite laths) Tso is associated with a low yield strength y ( 0.2 in first approximation) by the relationship Tso = y + C, where C and are constants. For heat 1 Tso varies linearly with o.2 in a broad range of changes in yield strength, and only with tempering at 400-600 , where special carbides begin to develop, is there a deviation from the linear relationship. Of course, these processes affect the boundary zones first, where conditions are favorable for the formation of carbide phase due to the geometric imperfection of the crystal structure. It can be assumed that the Tso peak is due to a change in the condition of the boundaries. This assumption is confirmed by the results of electron microscopic analysis: The condition of the matrix (dislocation density, size of fragments), causing an increase in the strength of the material 8, 9, remains unchanged up to Ttemper = 600 . This is confirmed by examination of the structure and by the variation of the yield strength with the tempering temperature. Thus, if no changes occurred in the boundary to increase embrittlement, then the variation of Tso with Ttemper would have the form of the dashed line. All theories explaining RTB 1, 2, 10 are based on the recognition of the dominant role of boundary effects. However, damage in the boundaries of prior austenite grains was not typical in our investigation (the fracture of heat 1 after quenching and after tempering was quasibrittle), and no effect of phosphorus was observed, which is evidently explained by the presence of molybdenum in the steel. It follows from the experimental data that the embrittlement of CMoV steels during tempering is affected mainly by carbide formation, more precisely the rebuilding of carbides and factors affecting these processes. The mechanism of this effect can be presented in general terms as follows. With increasing tempering temperatures cementite begins to coalesce, precipitating at 250-350 . In steels containing more effective carbide-forming elements than iron the formation of nuclei of special carbides begins around 350 , especially MC3 8. Precipitating evenly throughout the volume on disloca tions, low-angle boundaries of martensite laths, and high-angle boundaries of colonies of laths, they strengthen the matrix and weaken (embrittle) the boundaries. It is probable that the boundaries are weakened most with the maximum density of carbides coherent with the matrix or with their precipitation, i.e., at secondary hardening temperatures. Such elements as vanadium, promoting refining of carbides and thus increasing secondary hardening 8, increase the embrittlement. With precipitation and coalescence of carbide phase, occurring simultaneously with polygonization of dislocations, the boundaries become more perfect and the matrix is weakened. Both these effects lead to a drop of the ductile-brittle transition temperature, which is in fact observed. CONCLUSIONS1. Steels of the 15Kh3MFA type are susceptible to embrittlement, which reaches a peak at a given tempering temperature. The upper limit of the peak (Tso) coincides with the temperature at which the material begins to weaken.2. The height of the peak and the tempering temperature corresponding to it increase almost linearly when the vanadium concentration is raised from 0 to 0.55%, but they are independent of the phosphorus concentration within limits of 0.005-0.022%.3. Temper brittleness of the steel investigated depends on the change in the carbide phase (from cementite to special carbides) that occurs with retention of the dislocation arrays preferentially in the boundaries of fragments.LITERATURE CITED1.L. M. Utevskii, Temper Brittleness of Steels in Russian, Metallurgizdat,